Polymer Functionalized Graphene

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Polymer Functionalized Graphene

Polymer Chemistry Series Editor-­in-­chief:

Ben Zhong Tang, The Hong Kong University of Science and Technology, ­ Hong Kong, China

Series editors:

Alaa S. Abd-­El-­Aziz, University of Prince Edward Island, Canada Jianhua Dong, National Natural Science Foundation of China, China Jeremiah A. Johnson, Massachusetts Institute of Technology, USA Toshio Masuda, Shanghai University, China Christoph Weder, University of Fribourg, Switzerland

Titles in the series:

1: Renewable Resources for Functional Polymers and Biomaterials 2: Molecular Design and Applications of Photofunctional Polymers and Materials 3: Functional Polymers for Nanomedicine 4: Fundamentals of Controlled/Living Radical Polymerization 5: Healable Polymer Systems 6: Thiol-­X Chemistries in Polymer and Materials Science 7: Natural Rubber Materials: Volume 1: Blends and IPNs 8: Natural Rubber Materials: Volume 2: Composites and Nanocomposites 9: Conjugated Polymers: A Practical Guide to Synthesis 10: Polymeric Materials with Antimicrobial Activity: From Synthesis to Applications 11: Phosphorus-­based Polymers: From Synthesis to Applications 12: Poly(lactic acid) Science and Technology: Processing, Properties, Additives and Applications 13: Cationic Polymers in Regenerative Medicine 14: Electrospinning: Principles, Practice and Possibilities 15: Glycopolymer Code: Synthesis of Glycopolymers and their Applications 16: Hyperbranched Polymers: Macromolecules in-­between Deterministic Linear Chains and Dendrimer Structures 17: Polymer Photovoltaics: Materials, Physics, and Device Engineering 18: Electrical Memory Materials and Devices 19: Nitroxide Mediated Polymerization: From Fundamentals to Applications in Materials Science 20: Polymers for Personal Care Products and Cosmetics 21: Semiconducting Polymers: Controlled Synthesis and Microstructure 22: Bio-­inspired Polymers 23: Fluorinated Polymers: Volume 1: Synthesis, Properties, Processing and Simulation 24: Fluorinated Polymers: Volume 2: Applications 25: Miktoarm Star Polymers: From Basics of Branched Architecture to Synthesis, Self-­assembly and Applications

26: Mechanochemistry in Materials 27: Macromolecules Incorporating Transition Metals: Tackling Global Challenges 28: Molecularly Imprinted Polymers for Analytical Chemistry Applications 29: Photopolymerisation Initiating Systems 30: Click Polymerization 31: Organic Catalysis for Polymerisation 32: Synthetic Polymer Chemistry: Innovations and Outlook 33: Amphiphilic Polymer Co-­networks: Synthesis, Properties, Modelling and Applications 34: Redox Polymers for Energy and Nanomedicine 35: Polymer Functionalized Graphene

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Polymer Functionalized Graphene By

Arun Kumar Nandi

Indian Association for the Cultivation of Science, India Email: [email protected]

Polymer Chemistry Series No. 35 Print ISBN: 978-­1-­78801-­879-­1 PDF ISBN: 978-­1-­78801-­967-­5 EPUB ISBN: 978-­1-­78801-­968-­2 Print ISSN: 2044-­0790 Electronic ISSN: 2044-­0804 A catalogue record for this book is available from the British Library © Arun Kumar Nandi 2021 All rights reserved Apart from fair dealing for the purposes of research for non-­commercial purposes or for private study, criticism or review, as permitted under the Copyright, Designs and Patents Act 1988 and the Copyright and Related Rights Regulations 2003, this publication may not be reproduced, stored or transmitted, in any form or by any means, without the prior permission in writing of The Royal Society of Chemistry, or in the case of reproduction in accordance with the terms of licences issued by the Copyright Licensing Agency in the UK, or in accordance with the terms of the licences issued by the appropriate Reproduction Rights Organization outside the UK. Enquiries concerning reproduction outside the terms stated here should be sent to The Royal Society of Chemistry at the address printed on this page. Whilst this material has been produced with all due care, The Royal Society of ­ Chemistry cannot be held responsible or liable for its accuracy and completeness, nor for any consequences arising from any errors or the use of the information contained in this publication. The publication of advertisements does not constitute any endorsement by The Royal Society of Chemistry or Authors of any products advertised. The views and opinions advanced by contributors do not necessarily reflect those of The Royal Society of Chemistry which shall not be liable for any resulting loss or damage arising as a result of reliance upon this material. The Royal Society of Chemistry is a charity, registered in England and Wales, Number 207890, and a company incorporated in England by Royal Charter (Registered No. RC000524), registered office: Burlington House, Piccadilly, London W1J 0BA, UK, Telephone: +44 (0) 20 7437 8656. For further information see our website at www.rsc.org Printed in the United Kingdom by CPI Group (UK) Ltd, Croydon, CR0 4YY, UK

Preface Since the discovery of graphene in 2004 by Geim and Novoselov of University of Manchester, UK, it has drawn considerable attention from physicists, chemists, biologists and presently from technologists. Graphene is a two-­dimensional, sp2 hybridized flat crystalline form of carbon atoms in a hexagonal lattice structure making it a highly optically transparent and mechanically strong material with exciting electrical and heat conducting properties, leading to many applications in technology because of its high tensile strength (∼0.4 GPa), Young's modulus (500 GPa), much higher than steel, and elasticity more than rubber. Despite the important properties of graphene, it is difficult to process due to its strong π-­stacking interaction, and functionalization of graphene is utmost necessity for commercial use because of the reinforcing properties of graphene on polymers, imparting both the strong mechanical and good electrical properties of polymers. This is the reason for writing a book on polymer functionalized graphene (on invitation from the Royal Society of Chemistry), which is a field I have been working on for more than a decade, exploring different aspects polymer functionalized graphene. The book consists of twelve chapters, the first one gives an introduction to graphene, including its different preparation procedures, properties, formation of graphene oxide (GO), reduced graphene oxide (rGO) and graphene quantum dot (GQD). The characterization of the above materials using microscopic X-­ray and different spectroscopic techniques, including optoelectronic properties of the materials, are discussed. Here the necessity of functionalization, specifically the importance of polymer functionalization is also embodied. In the second chapter, the various synthetic protocols of covalent grafting of polymers from the graphene surface, e.g. esterification, amide formation, click chemistry, free radical polymerization, atom transfer   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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radical polymerization (ATRP), reversible addition−fragmentation chain-­ transfer (RAFT) polymerization, etc. are discussed from both the grafting to and grafting from approaches. The characterization of these covalently functionalized polymer functionalized graphenes (PFGs) is also briefly included here. In the third chapter, the necessity of non-­covalent functionalization and different methods of noncovalent functionalization using dispersion, π-­stacking, electrostatic, coordination and hydrogen bonding interactions for producing PFGs are discussed. A comparison between covalent and noncovalent functionalized PFGs is also made in respect of their properties, emphasizing the specific utility of noncovalently functionalized PFGs. In the fourth chapter, the morphology, structure, physical and thermal properties of different PFGs are discussed. The influence of graphene on the change of morphology, structure, thermal stability, glass transition temperature, melting temperature, crystallinity, and polymorphic structure of the polymers of the PFGs are delineated with the physico-­chemical background behind it. The fifth chapter discusses how polymer functionalization of GO influences the UV–vis, Raman and fluorescence spectra of different PFGs, including the applications in different optical sensors for explosives, toxic ions and biomolecules. The sixth chapter discusses the improvement of the mechanical properties of PFGs and their composites with other polymers because graphene causes a significant improvement in storage and loss modulus, tensile stress, tensile strain, Young's modulus, etc. due to its high aspect ratio and good reinforcing properties. A thorough analysis of the results in different composite systems using theoretical models is made. In Chapter 7 electronic and ionic conductivity, and current–voltage (I–V) behavior of PFGs and their composites are discussed for both nonconducting and conducting polymer systems. The results are analysed with p–n junction formation, p-­and n-­t ype doping, negative differential resistance, etc. In Chapter 8,the dielectric permittivity and dielectric loss properties of PFGs are discussed where graphene causes an increase in the dielectric constant and decrease in the dielectric loss in a polymer matrix due to formation of a microcapacitor network. In addition, sometimes low dielectric constant values of PFGs/polymer composites are necessary to a lower value of 2, particularly for the electronics industry, and it is also noticed in some PFGs. Some probable explanation of increases and decreases in the dielectric properties of these composites are discussed. The perspectives of these dielectric systems for applications in power industry and semiconductor technology is also highlighted. Chapter 9 deals with application of PFGs with conducting polymers like polyaniline, polypyrrole and polythiophene in photovoltaics are discussed. Both covalent and noncovalent functionalized polymer–graphene nanocomposites suitable for photovoltaic energy generation are discussed, delineating their use in fabrication of bulk heterojunction, dye sensitized and perovskite solar cells. In Chapter 10, the use of PFGs in fuel cells is discussed for hydrogen and methanol fuel cells as examples. The PFGs are used as proton exchange membrane, anion exchange membrane and also as an electrocatalyst and the

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fuel cell parameters show a significant improvement from those of pristine polymers. In Chapter 11 the use of PFGs in a solid state battery and supercapacitors is discussed; in the former, the use of PFGs as anode, cathode and electrolyte are discussed with some examples. Solid polymer electrolyte materials with reduced crystallinity and higher ionic conductivity are very much necessary and here their development is discussed using PFGs. The PFGs are highly used in flexible supercapacitors because of the large specific surface area, high mechanical stability and good conducting property of graphene or reduced graphene oxide. Here polymer functionalized GO both by covalent and noncovalent ways are discussed for their use as electrode materials of supercapacitors. Both symmetric and asymmetric super capacitor devices fabricated with PFG electrodes are also discussed with an aim to improve the specific capacitance, power density of the device and also for their long term stability. In Chapter 12, the PFGs produced with biocompatible polymers are found to be highly useful in ultrasensitive biosensors, drug delivery, gene delivery, cell imaging, smart implants, wound managements, etc. In each field of biotechnological applications of PFGs they are discussed with some specific examples elaborating the general principles involved along with their technological overview. So, in the first seven chapters of the book, the synthesis, characterization, physical, mechanical, optical and electronic properties of PFGs are embodied while the last five chapters embody the technological applications of PFGs, e.g. dielectric material, photovoltaics, fuel cells, solid state batteries, supercapacitors, and diverse biomedical and biotechnological applications. The author envisages an enormous growth of polymer functionalized graphene in technological fields, specifically, a huge burst will occur in applications for energy and different biomedical and biotechnological devices. Arun Kumar Nandi IACS, Kolkata-­32, India

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Acknowledgement I am grateful to my parents, teachers, students, colleagues, friends and collaborators of IACS and other universities who helped me a lot in the genesis of the work and also providing literature during COVID-­19 lockdown period. I show my gratitude to my mentors, specially to Professor B. M. Mondal who encouraged me to write the book. I sincerely thank my wife (Bithika), son (Amrit), and daughter (Debasmita) who give me inspiration all through the writing of the book. I also owe to my brothers and sisters and villagers of my native village “Belboni, Bankura, WB” whose love and inspiration helped make the book be a reality. Here it is necessary to specially mention few names of students; Drs Ramakanta Layek, Aniruddha Kundu, Atanu Kuila, Nabasmita Matity, Shreyam Chatterjee, Nirmal Maity, Amit Mondal, Parimal Routh, Arnab Shit and Debashis Mondal whose pioneering work in different fields of polymer functionalized graphene make me enough competent to gather knowledge for writing the book. I acknowledge Professors S Malik, D. P. Chatterjee,  P. Maity, T. Jana, T. K. Mondal, S. Ghosh, A. Banerjee, N. Jana, D. Chattopadhyay,  N. Mishra and A. Ghosh for their help and encouragement. I also thank my present students, S. Mondal, A. Panja, M. Pakhira, S. Hazra and U. Halder for their help in different ways. I am very much indebted to the help extended by Mr Amit Chakraborty for typing and Mr Gopal Manna who drew all the figures and graphics taking extra pain outside their routine work. I also acknowledge the help of our staff members, Siddartha, Champa, Subhashish and Mahadev for their help in various ways. Finally, I acknowledge the Director of IACS for providing me with infrastructural and library facilities, and CSIR, New Delhi for providing me with an Emeritus Scientist position to work at IACS. Professor Arun Kumar Nandi, FASc Polymer Science Unit, School of Materials Science Indian Association for the Cultivation of Science Jadavpur, Kolkata-­700 032, India   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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Dedicated To ‘MA’

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Contents Chapter 1

I ntroduction  1.1 Introduction  1.2 Short History of Graphene  1.3 Synthesis of Graphene  1.4 Graphene Oxide  1.4.1 Synthesis of GO  1.5 Reduced Graphene Oxide (rGO)  1.5.1 Synthesis of rGO  1.6 Characterization of Graphene/Graphene Oxide  1.6.1 Microscopy  1.6.2 Spectroscopy  1.7 Necessity for the Functionalization of Graphene  1.7.1 Necessity of Polymer Functionalization  1.8 Applications  1.8.1 Applications of Polymer Functionalized Graphene  1.9 Scope  References 

Chapter 2 Covalent Functionalization of Polymers  2.1 Covalent Functionalization  2.2 ‘Grafting to’ Method  2.2.1 Esterification Reaction  2.2.2 Amidation Reaction  2.2.3 Click Chemistry  2.2.4 Nitrene Chemistry  2.2.5 Radical Addition  2.2.6 Other Methods   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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1 1 1 3 5 6 6 7 9 9 11 16 17 18 18 18 19 24 24 24 25 29 34 36 40 45

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Chapter 3

2.3 ‘Grafting from’ Method  2.3.1 Atom Transfer Radical Polymerization (ATRP)  2.3.2 Reversible Addition Fragmentation Chain Transfer (RAFT) Polymerization  2.4 ‘Grafting to’ Versus ‘Grafting from’ Technique  2.5 Scope  References 

58 65 66 66

 oncovalent Polymer Functionalization of Graphene  N 3.1 Introduction  3.2 π-­stacking Interaction  3.3 H-­bonding Interactions  3.4 Surfactant Induced Functionalization  3.5 Miscellaneous Nonbonding Interactions  3.6 Mixed Noncovalent and Covalent Functionalization  3.7 Scope  References 

72 72 72 80 84 87 90 92 92

Chapter 4 P  hysical Properties of Polymer Functionalized Graphene  4.1 Morphology  4.1.1 Transmission and Scanning Electron Microscopy  4.1.2 Atomic Force Microscopy (AFM)  4.2 Structural Study  4.2.1 Fourier Transformed Infrared Spectroscopy (FTIR)  4.2.2 Raman Spectroscopy  4.2.3 X-­ray Photoelectron Spectroscopy (XPS)  4.2.4 Wide Angle X-­ray Scattering (WAXS)  4.3 Thermal Properties  4.3.1 Thermogravimetric Analysis (TGA)  4.3.2 Differential Scanning Calorimetry (DSC)  4.4 Conclusion  References  Chapter 5 O  ptical Properties of Polymer Functionalized Graphene: Application as Optical Sensor  5.1 Introduction  5.2 UV–Vis Spectra  5.3 Photoluminescence (PL) Spectra  5.3.1 Fluorescence in Polymer Functionalized Graphene  5.3.2 Fluorescence Quenching  5.3.3 Fluorescence Properties of PDMAEMA Grafted rGO (RGP): pH Dependent Doping 

47 48

95 95 95 102 105 105 109 113 116 121 121 125 129 129

133 133 133 136 137 144 145

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5.3.4 Fluorescence Properties of GO-­g-­poly (ϵ-­caprolactone) (PCL)-­b-­poly(dimethyl Aminoethyl Methacrylate) (GPCLD): LCST, Sensing and pH Dependent Doping  5.3.5 Fluorescent Amylose-­functionalized Graphene: Chiral Detection  5.3.6 β-­Cyclodextrin Functionalized Graphene: Fluorescent Detection of Cholesterol  5.3.7 Fluorescent Block Copolymer-­functionalized Graphene Oxide: Efficient Temperature Sensing  5.4 Scope  References 

Chapter 6 M  echanical Properties of Polymer Functionalized Graphene  6.1 Introduction  6.2 Dynamic Mechanical Properties  6.2.1 Covalently Functionalized Graphene Nanocomposites  6.2.2 Noncovalently Functionalized Graphene Nanocomposites  6.2.3 Functionalized Graphene/Polystyrene Composites  6.3 Mechanical Properties  6.3.1 Covalently Functionalized Graphene Nanocomposites  6.3.2 Noncovalently Functionalized Graphene Nanocomposites  6.4 Conclusion  References  Chapter 7 E  lectronic Properties of Polymer Functionalized Graphene  7.1 Introduction  7.2 Conductivity  7.2.1 Dc Conductivity  7.2.2 Ac Conductivity  7.2.3 Conclusion  7.3 Ionic Conductivity  7.3.1 Proton Conductivity  7.3.2 Hydroxide Ion Conductivity  7.3.3 Conclusion  7.4 Current–Voltage (I–V) Properties  7.4.1 Covalently Functionalized Systems  7.4.2 Noncovalent Functionalized Systems  7.5 Conclusion and Future Perspectives  References 

148 154 157 159 161 161 164 164 165 165 170 174 175 175 185 193 194 197 197 198 198 208 211 211 212 215 217 217 217 224 229 229

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Chapter 8 P  olymer Functionalized Graphene as Dielectric Material  8.1 Introduction  8.2 Covalent Functionalized Graphene/Polymer Systems  8.2.1 Systems with Increased Dielectric Properties  8.2.2 Systems with Decreased Dielectric Properties  8.3 Noncovalent Functionalized Graphene/ Polymer Systems  8.3.1 Systems with Increased Dielectric Properties  8.3.2 Systems with Low Dielectric Properties  8.4 Conclusion and Perspective  References  Chapter 9 A  pplications of Polymer Functionalized Graphene in Energy Harvesting: Photovoltaics  9.1 Introduction  9.2 Bulk Heterojunction (BHJ) Solar Cells  9.2.1 Graphene/PEDOT:PSS/(P3HT-­PCBM)/ZnO Based BHJ Solar Cell  9.2.2 Graphene/PEDOT/CuPc/C60/BCP Based BHJ Solar Cell  9.2.3 Graphene Oxide/PEDOT:PSS Based BHJ Solar Cell  9.2.4 Graphene Nanoflakes/(PCDTBT/PC 71 BM) Based BHJ Solar Cell  9.2.5 rGO-­PEDOT:PSS/PANI-­Ru Based BHJ Solar Cell  9.3 Dye Sensitized Solar Cells  9.3.1 Replacement of TiO2 Active Layer in the Photoelectrode with an rGO Grafted PANI System  9.3.2 Replacement of the TiO2 Active Layer with Poly(Hydroxyethyl Thiophene) Grafted rGO  9.3.3 Replacement of the TiO2 Active Layer with a GQD/PT-­g-­P(MeO2MA-­co-­DMAEMA) Hybrid  9.3.4 Replacement of the TiO2 Active Layer with a Graphene Quantum Dot/PANI Hybrid  9.3.5 Replacement of TiO2 Active Layer with Graphene/Polymer Hybrid Xerogels  9.4 Replacement of the Pt Counter Electrode with PFG  9.5 Improving Electrolyte Performance with PFG 

Contents

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9.6 Perovskite Solar Cell  9.6.1 Poly[(5,6-­difluoro-­2,1,3-­benzothiadiazol-­4, 7-­diyl)-­alt-­(3,3‴-­di(2-­octyldo decyl)-­2,2;5, 2;5,2″-­quaterthiophen-­5,5‴-­diyl)] as Hole Transporting Layer (HTL)  9.6.2 Graphene–AgNWs–Polycarbonate (MG-­A-­P) Film as Electron Transporting Layer  9.6.3 TFSA-­doped Graphene PDMS as Hole Transporting Flexible Electrode  9.7 Controlling Grain and Crystal Size of Perovskites Using Polymer Additive and Enhancing PCE and Stability of the Cell  9.8 Conclusion  References 

Chapter 10 A  pplications of Polymer Functionalized Graphene in Energy Harvesting: Fuel Cells  10.1 Introduction  10.2 Polymer Functionalized Graphene in Fuel Cells  10.2.1 Hydrogen Fuel Cell  10.2.2 Polymer Functionalized Graphene as Anion Exchange Membrane (AEM)  10.2.3 Methanol Fuel Cell  10.3 Conclusion  References  Chapter 11 P  olymer Functionalized Graphene in Energy Storage Devices  11.1 Introduction  11.2 Solid State Battery  11.2.1 PFG as Anode in a Solid State Battery  11.2.2 PFGs as Cathode in Solid State Battery  11.2.3 PFGs as Electrolyte in a Solid State Battery  11.3 Supercapacitors  11.3.1 Graphene as a Good Supercapacitor Material  11.3.2 Polymer Functionalized Graphene: As an Efficient Supercapacitor Material  11.4 Conclusion  References  Chapter 12 P  olymer Functionalized Graphene in Biomedical and Bio-­technological Applications  12.1 Introduction  12.2 Polymer Functionalized Graphene as Biosensors  12.2.1 Dopamine-­functionalized Polyethylene Glycol and 2,5-­thiophenediylbisboronic Acid Conjugated Graphene as Fluorometric Biosensors of Glucose and E. coli 

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283 285 286 288 289 289 292 292 295 295 308 312 318 318 322 322 324 324 328 333 337 338 341 353 354 357 357 358

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12.2.2 Graphene Quantum Dots and Pyrene-­ functionalized Molecular Beacon Probes for Fluorimetric Sensing of MicroRNA  12.2.3 PEGMA, Oligonucleotides and GO Nanoassembly for Fluorimetric Detection of DNA, miR-­10b, Thrombin and Adenosine  12.2.4 GO Based Molecular Imprinted Polymer (GO-­MIP) for Amperometric Detection of Total Cholesterol  12.2.5 GO-­copolymer Hybrid for Amperometric Detection of Dengue Virus  12.2.6 GO-­PANI/Ag or Au NP Hybrid for Amperometric Detection of Vitamin C  12.3 Polymer Functionalized Graphene for Application in Drug Delivery  12.3.1 PNIPAM-­grafted GO for Delivery of Both Hydrophilic and Hydrophobic Drugs  12.3.2 Starch Functionalized Graphene for pH Sensitive and Starch-­mediated Drug Delivery  12.3.3 GO Conjugated Chitosan for the In Vitro and In Vivo Co-­delivery of Anti-­cancer Drugs  12.3.4 Polyurethane Grafted Sulphonated Graphene as Drug Delivery Vehicle  12.3.5 Poly(Vinyl Pyrrolidone)-­functionalized GO as a Nanocarrier for Dual Drug Delivery  12.4 Polymer Functionalized Graphene for Application in Gene Delivery  12.4.1 GO Grafted with Positively Charged PEI for Transfection of Plasmid DNA  12.4.2 Injectable PEI-­functionalized GO Hydrogel Based Angiogenic Gene Delivery System for Vasculogenesis and Cardiac Repair  12.4.3 Polyethylenimine and Polyethylene Glycol Dual-­functionalized GO for High-­efficiency Delivery of DNA and siRNA  12.4.4 Peptide Functionalized GO Nanocarrier for Gene Delivery Applications  12.4.5 GO–Chitosan Nanocomposites for Intracellular Delivery of Immunostimulatory CpG Oligodeoxynucleotides  12.5 Polymer Functionalized Graphene for Application in Cell Imaging  12.5.1 RGO Nanoribbons Functionalized with Polyethylene Glycol (rGONR–PEG) for Cell Imaging 

359 362 363 364 366 367 367 370 371 373 374 376 376 377 378 381 382 384 385

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12.5.2 GQDs Functionalized with Hyaluronic Acid for Cell Imaging  12.5.3 Magnetic GO Functionalized with Cyclodextrin–Hyaluronic Acid Polymer for Cancer Cell Imaging  12.6 Polymer Functionalized Graphene in Tissue Engineering  12.6.1 GO Functionalized with Polypeptide of l-­lysine for Cardiac Tissue Engineering  12.6.2 Polymer Functionalized Graphene for Neural Tissue Engineering  12.7 Polymer Functionalized Graphene in Body Implants  12.7.1 GO/Alginic Acid (AA, a Natural Polymer)/ a Bioceramic (TCP) Composite for Bone Implant  12.7.2 GO/Polycaprolactone/Hydroxyapatite Based Bioactive Coating on Ti Alloy for Bone Implant  12.7.3 Amino Functionalized Graphene/ Poly(Methyl Methacrylate-­co-­styrene) Copolymer for Effective Bone Cement Implant  12.7.4 GO-­Polyetheretherketone for Orthopedic Implant  12.8 Polymer Functionalized Graphene for Wound Healing Applications  12.8.1 Ag/Graphene-­polymer Hydrogel for Antibacterial and Wound Healing Application  12.8.2 rGO-­poly(Diallyldimethylammonium Chloride)/Ag/AgCl Hybrid Material for Burn Wound Healing  12.8.3 Polydopamine-­rGO in Mussel-­inspired Electroactive and Antioxidative Scaffolds for Enhancing Skin Wound Healing  12.8.4 rGO-­isabgol Nanocomposite Dressings for Enhanced Vascularization and Accelerated Wound Healing  12.8.5 TRGO-­polydopanine-­boronic Acid System for Diabetic Wound Healing  12.9 Conclusion  References 

Subject Index

386 388 389 389 392 402 402 404

406 407 410 411 413 414 416 418 419 420 426

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Chapter 1

Introduction 1.1  Introduction After the invention of X-­ray crystallography, the layered structure of graphite was unveiled and was found to be a deck of weakly bonded graphene planes. Graphene is a two-­dimensional, sp2 hybridized flat crystalline form of carbon atoms in a hexagonal lattice structure (Figure 1.1).1–3 This highly optically transparent and mechanically strong material has exciting electrical properties and heat conductivity, leading to many applications in technology. It has a C–C bond length of 0.142 nm and an atomic thickness of 0.345 nm. Its tensile strength of ∼0.4 GPa is much higher than steel, its Young's modulus is 500 GPa, and it is more elastic than rubber.4 Due to these incredible properties it is much studied in chemistry, physics, materials science, engineering and also in biology.

1.2  Short History of Graphene Sporadic attempts at graphene synthesis started in 1859,5 but pre-­2004 results were experimental, where they obtained ultra-­thin graphitic films, but did not report any of graphene's characteristic properties. Earlier attempts at producing graphene were concentrated on a chemical method of exfoliation. Bulk graphite was first intercalated6 to separate layers by intervening atoms or molecules. The insertion of large molecules between atomic planes causes greater separation so that it results in isolated graphene layers into a 3D matrix. However, due its uncontrollable nature, the graphitic

  Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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Figure 1.1  The  crystal structure of graphene – carbon atoms arranged in a hon-

eycomb lattice. Figure By AlexanderAlUS, Reproduced from https:// en.wikipedia.org/wiki/Graphene#/media/File:Graphen.jpg under the terms of the CC BY-­SA 3.0 license, https://creativecommons.org/ licenses/by-­sa/3.0/.

muck found very low interest. In another method, few-­layer graphene was grown epitaxially by chemical vapour deposition (CVD) of hydrocarbons on metal substrates.7,8 Also in other techniques, SiC on thermal decomposition yielded graphene films that were characterized using surface science techniques.9,10 In 2004, Novoselov et al. first isolated graphene in large quantity using a simple ‘Scotch tape method’.1,2 This micromechanical cleavage (Scotch tape) method was very simple and effective for growing graphene science very quickly. This technique does not require sophisticated equipment, hence it helped to grow graphene science enormously. For their new Scotch tape method of synthesizing graphene, correct characterization and delineating the beautiful physics of graphene, Geim and Novoselov of University of Manchester, UK obtained the Nobel Prize in Physics in the year 2010. Another physical method to prepare graphene is the thermal expansion of graphite followed by ultra-­sonication in an aqueous medium yielding dispersion of graphene sheets11,12 and this has also gained popularity to produce graphene. Then chemists become interested in synthesizing graphene by oxidizing graphite followed by exfoliation on sonication and the produced graphene oxide is then reduced chemically to get reduced graphene oxide in plentiful amounts. The preparation methods of graphene, graphene oxide, reduced graphene oxide and graphene quantum dots are described in detail below.

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1.3  Synthesis of Graphene The different methods of producing graphene from graphite are presented in Scheme 1.1. It can be produced in a ‘top–down’ method to obtain a single-­layer graphene from graphite crystals, or using a ‘bottom–up’ method to form graphene on self-­assembling small molecules used to create graphene, respectively. Among the top–down methods the mechanical peeling by the ‘Scotch tape technique’, forms a single-­layer graphene material. This is made by repeated peeling of small mesas of greatly oriented pyrolytic graphite. It is the first method to produce large quantities of good quality of graphene, and is the original method of its invention.1,2 The films are two-­dimensional semimetal having small overlap between conductance and valence bands, exhibiting a ambipolar electric field effect generating holes and electrons in concentrations up to 1013 cm−2 with mobilities of about 10 000 cm2 V−1 S−1 on application of gate voltage. Another top–down technique is the liquid-­phase exfoliation where delamination is done using exfoliating agents (usually solvents) to prevent the overlap between adjacent layers by making colloidal dispersions. For example, graphene powder mixed with N-­methyl pyrrolidone (NMP) is sonicated vigorously producing a colloidal dispersion of graphene and removing the colloidal aggregates by centrifugation.13 This is possible because the required energy to exfoliate graphene is coming from the graphene–solvent interaction, particularly for solvents whose surface energies match with that of graphene. N,N′ dimethyl formamide (DMF) has also the similar property to exfoliate graphene and the yield of graphene could be up to 12% of the graphite used for exfoliation. Then a stabilizer is used to stop restacking of the graphene layers and thus produce stable graphene sheets.

Scheme 1.1  Schematic  illustration of different methods to obtain graphene from graphite Reproduced from ref. 24 with permission from Elsevier, Copyright 2015.

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Chapter 1

In a bottom-­up technique of producing graphene, chemical synthesis is used to produce small-­sized graphene or graphene nanoribbon (GNRs). In a microwave plasma reactor substrate-­free gas-­phase synthesis of graphene is well reported.14 Also arc discharge synthesis15 of multi-­layered graphene is reported. GNRs are produced from the reaction between high aromatic rings and polycyclic aromatic hydrocarbons. However, the most current and trending method to create graphene from the bottom–up approach is via chemical vapour deposition (CVD), which is well established, providing the material in sufficient quantity and requires an easy laboratory set-­up. In the CVD process, reactive gases (methane, etc.) at a regulated rate are passed through a gas-­mixing unit mixing the gases uniformly before feeding into the reactor. In the reactor a chemical reaction occurs and the solid products become deposited on the metal substrates. The reactor has a heating arrangement to maintain the high temperatures required for the reaction. The source of C is mostly CH4, which is catalytically dehydrogenated on the Cu surface, then a solid solution of C is formed near the metal surface, resulting in graphene. This route facilitates surface migration and monolayer graphene growth. A schematic model of the working process is shown in Figure 1.2.16 The by-­products of reaction with the unused gases get removed by the gas outlet system. Graphene is grown on the surface of metals by decomposition of hydrocarbons or by segregation of carbon.17 To improve graphene formation and to obtain a smooth graphene surface, treatment of the metal substrate before beginning the CVD process is generally done. Wet chemical treatments are made for metal films to avoid oxide reduction by soaking in acetic acid, etc. Then, a metal substrate, is heated under low vacuum in

Figure 1.2  Schematic  illustration of the chemical vapour deposition process: (1)

Diffusion of reactant precursors through the boundary layer. (2) Adsorption of reactants onto the substrate surface. (3) Onset of a chemical reaction at the substrate surface. (4) Desorption of the adsorbed product from the surface. (5) diffusion of by-­products through the boundary layer. Reproduced from https://commons.wikimedia.org/wiki/File: Sequence_during_CVD_%28en%29.svg under the terms of the CC BY-­ SA 3.0 license, https://creativecommons.org/licenses/by-­sa/3.0/deed.en.

Introduction

5

the CVD process. Good quality graphene can be grown on Cu surfaces by the CVD process;18 however on Ni surfaces some roughness of graphene surface is observed.19,20 In Ni the solubility of C is high, making it difficult to suppress C precipitation and the thickness of produced graphene films vary from single to ten layers. However, use of Cu foils in CVD produces large-­ area graphene films containing 95% monolayer graphene due to very low solubility of carbon in Cu.18 Li et al. has shown that growth of graphene on a Cu surface is self-­limiting, i.e. when the Cu surface is entirely coated with graphene, the growth becomes terminated, producing almost 100% monolayer graphene.21 Graphene produced from the CVD process is of high quality showing carrier mobility in the range 2000–4000 cm2 V−1 s−1.

1.4  Graphene Oxide Graphene oxide (GO) is different from a graphene sheet that contains only sp2 hybridized carbon atoms in that the GO sheet contains a hexagonal carbon network having sp2-­hybridized carbon along with sp3-­hybridized carbons containing oxygen functional groups (Figure 1.3).22–24 In GO, the oxygen functional groups are hydroxyl, carboxyl and epoxy, which are linked to the sp3 hybridized carbon atoms. Hence the conjugated network of sp2-­hybridized graphene becomes disrupted decreasing the extended π-­conjugation of graphene. The sp3 carbon clusters are randomly located somewhat top and bottom of the graphene plane.25 Transmission electron microscopy (TEM), atomic force microscopy (AFM) and scanning tunnelling microscopy (STM) yields the length, breadth and thickness of single-­layer GO of around 1 nm, hence the number of layers present in GO is determined. It lacks the ordered lattice of graphene for the presence of oxygenated functional groups.23 The oxygen atoms remain linked to the carbon atoms of graphene randomly converting planar sp2 carbon bonds of graphene to tetrahedral sp3 bonds. The oxygen functionalities in GO are identified using different spectroscopic methods, like Fourier transform infrared (FT-­IR),26,27 nuclear magnetic resonance,28,29 Raman26,30 and X-­ray

Figure 1.3  Conversion  of graphene to GO using Hummer's method. Reproduced from ref. 24 with permission from Elsevier, Copyright 2015.

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6 31,32

photoelectron (XPS) spectroscopy. The formation of GO is made by oxidizing graphite with KMnO4/NaNO3 in sulphuric acid medium using Hummer's method33 (Figure 1.3) where the structure of GO containing different oxygen functionalities are clearly presented.

1.4.1  Synthesis of GO GO is an aromatic framework of graphene consisting of alcohol, carbonyl, epoxide and carboxylic groups. The introduction of the above functional groups in the graphene lattice causes an increase of interlayer spacing of graphene from 0.34 nm for graphite to higher than 0.63 nm in GO.35 In 1859 Brodie was the first to synthesize GO by adding a portion of potassium chlorate to a slurry of graphite in fuming nitric acid.5 Staudenmaier advanced this protocol by taking a mixture of fuming nitric acid and concentrated sulphuric acid and finally adding chlorate in multiple aliquots during the reaction and this change produced highly oxidized GO.36 Hummer, in 1958, improved the oxidation method for exfoliation of graphene by oxidizing the graphite with NaNO3 and KMnO4 in concentrated H2SO4 and generally this method is commonly used34 (Figure 1.3).Graphite flakes are commonly used for oxidation purposes and the oxidized product in water is vigorously sonicated to obtain the GO sheets as thin films. In another method the preparation of GO nanoribbons (GONRs) is reported by treating multiwalled carbon nanotubes with KMnO4 and a concentration of H2SO4/H3PO4 mixture without using NaNO3 and this reaction has produced GONRs with intact basal planes of graphene with fewer defects than observed from Hummer's method.37,38 Very recently GO was synthesized from graphite powder in considerable amounts using a modified Hummer's method.39 Here graphite powder, and NaNO3 are mixed together then a concentration of H2SO4 is mixed at constant stirring. After one hour KMnO4 is gradually added to the solution at the temperature 4 the pendent –COOH groups become deprotonated, decreasing the PL intensity significantly. On esterification of the anchored –COOH group of GO, the PL quenching at the pH range 4–6 has been prevented suggesting that the electron of carboxylate ion is accountable to the PL quenching. Above pH 7 the pendent –OH group of GO becomes deprotonated causing further quenching due to a similar nonradiative recombination. Ajayan and his co-­workers have proposed that the fluorescence properties of GO originate from the quasi molecular fluorophores like in the polycyclic aromatic compounds due to the strong electronic coupling between oxygen atoms of carboxylic groups and the nearby carbon atom of graphene using the semi-­empirical quantum mechanical calculation.77 They noticed strongly pH-­dependent fluorescence of GO and it has been suggested that the PL properties of GO above pH 8 are due to the electronically excited carboxylate ion of GO. However, the fluorescence in GO quantum dots (GQDs) is pH dependent in contrast to that of GO, i.e. deprotonation causes increased emission, vis-­a-­vis protonation causes quenching. In the GQDs the emission is thought to originate from the zigzag sites with a carbene like triplet (σ1π1) ground state.77

1.7  Necessity for the Functionalization of Graphene Graphene has potential applications in different fields of material science for developing sensors, supercapacitors, nanocomposites, hydrogen storage, optoelectronic and biotechnological devices.78–83 The burgeoning applications of graphene lies in its function based on its fundamental properties, e.g. its excellent mechanical,4 thermal,17 electrical84 transport,85 thermoelectric86 and gas barrier properties.87 The strong cohesive force between the graphene layers in graphite causes difficulties in getting exfoliated graphene sheets. To reduce the cohesive force between the graphene sheets and also to impart specific interaction with the host matrix functionalization of graphene is of utmost necessity. On oxidation graphite produces graphite oxide (GO) containing functional groups like –COOH, –OH, epoxide, which are hydrophilic in nature and promote the intercalation of water molecules into the gallery of graphene sheets. This causes easy detachment from each other by sonication, yielding highly dispersible GO sheets in water medium.88–91 These exfoliated sheets are generally used for different applications and for targeted applications they are further functionalized. Substantial progress in attaching small organic molecules on graphene surface has been made, exploiting the rich chemistry of carboxyl, hydroxyl and epoxy groups of GO. Covalent attachment of organic moieties on the GO surface with azobenzene, porphyrins and phthalocyanines, covalently attached to obtain stimulating optoelectronic

Introduction

17 92–96

properties. For the extraordinary light extinction coefficients in the visible region, porphyrins and phthalocyanines are attached to GO for its application as light harvesting antennas for energy from sunlight. Porphyrins are attached via amide bond formation using the reaction between amine functionalized porphyrins and carboxylic acid groups of GO.94 By reacting GO with octadecylamine (ODA) organophilic graphene is produced and it is then reduced with hydroquinone using a reflux process. However, concurrent functionalization and reduction of GO is also made by refluxing GO and octadecylamine in the absence of any reducing agents.97–99 The presence of a long octadecyl chain makes the hydrophilic GO hydrophobic in nature and GO–ODA transforms into an electrically conducting material due to its reduction. The major drawback of small molecular functionalization is that it faces difficulties in the formation of stable dispersion100–105 Also the stress transfer is lower in the small molecular functionalized system compared to that of polymer functionalized graphene as the number of functional groups present in small molecules is lower causing weaker interfacial interaction. In order to facilitate good dispersion and good interfacial interaction polymer functionalization of graphene surface is, therefore, very much necessary.

1.7.1  Necessity of Polymer Functionalization The fascinating properties of graphene have promoted a great deal of interest in graphene research. To exploit its intriguing properties, inclusion of graphene sheets within a polymer matrix having good dispersion with fine interface control is of supreme necessity. The GO is electrically insulating but becomes conducting when it is reduced to produce reduced graphene oxide (rGO). However, the GO sheets in organic solvent and rGO sheets in both aqueous and organic solvent undergo aggregation due to the high cohesive interaction making them hard to disperse. This lack of uniform dispersion causes difficulties in its utilization in commercial applications because of the weak interfacial interaction of GO/rGO with the host polymer matrix cast from these solvents.56,106 So for a stable uniform dispersion of graphene sheets in a polymer matrix polymer functionalization is needed, which reduces the cohesive force significantly for its large size and also increases the interfacial interaction between the filler and matrix due to the presence of a large number of functional groups present in the polymer chain. The interaction between polymer functionalized graphene and host polymer results in a condensed interface facilitating easy stress transfer. For strong specific interaction of graphene with the host polymer matrix highly soluble graphene sheets are made by covalent grafting with the synthetic polymers, and also by making noncovalent attachment of the polymers. Thus, the functionalization of the graphene surface with functional polymers permits the fabrication of high performance conducting composite materials, and a great deal of interest has been shown in polymer functionalized graphene research because of their wide range of applications, from sensing, dielectrics, energy, etc., and also in the vast world of biotechnology.

18

Chapter 1

1.8  Applications Due to its high mechanical strength, elasticity and conductivity, graphene affords many uses ranging from electronics to batteries. The global market of graphene products is gradually increasing due to its unique properties. The virtual transparency makes it useful for touch screens and the high conductivity makes it great for supercapacitor usage.53a Graphene's nano dimensional size allows for multiple uses in electronics, which may potentially outperform current silicon-­based products. Graphene has the highest electronic mobility at room temperature making it greater switching speed in electronics. Due to its light weight its high surface area with better chemical tolerance and high energy density, graphene is used in solid state batteries; now, incorporation of graphene has improved the durability and charge capacity of lithium-­ion battery anodes.107–109 Graphene enhances the performance of lithium-­ion batteries where the process of inserting and extracting Li+ ions is faster, easier and can be reversed. Graphene's characteristics make this one of the most promising of materials to pique interest this century. Companies and university research laboratories have spent and incredible amount of time and money in making this material even more adaptable for everyday use: from electrical devices to batteries, contact lenses to structural support for the preservation of antiquities.

1.8.1  Applications of Polymer Functionalized Graphene An enormous variety of research is now ongoing on polymer functionalized graphene (PFG), enhancing the properties of the polymer and also graphene. On the one hand, the physical, mechanical and electrical properties of the polymers are highly enhanced. On the other hand, the optoelectronic properties of graphene, cell viability of graphene, doping of graphene, etc.110 are highly improved. Utilizing these properties, scientists and technologists have found uses for graphene polymer hybrids in chemical and biological sensing, photovoltaic devices, supercapacitors, batteries, dielectric materials, drug/gene delivery vehicles, etc. This book sheds light on the synthesis, properties and application of polymer functionalized graphene, enlightening both fundamental and technological development.

1.9  Scope So it is apparent from the above discussions that polymer functionalized graphene (PFG) not only improves the properties of graphene itself but also enhances the properties of polymers significantly. In a word a synergistic improvement of the properties of PFG is expected. This would certainly be an area of much curiosity both for academic and technological persons as it is appropriate for practical applications aiming at the betterment of our civilization. Apart from the engineering applications, the high demand for energy and crucial problems of biotechnology can be tackled successfully with PFG.

Introduction

19

For any useful application of PFG the basic science involved behind it should be known very well. In this book, an attempt has been made to highlight the fundamentals of polymer functionalized graphene, discussing its different synthetic details in covalent and noncovalent ways, characterization, properties and applications in various fields in the following chapters.

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Introduction

23

92. L. Xu, W. J. Yang, K. Neoh, E. Kang and G. D. Fu, Macromolecules, 2010, 43, 8336. 93. L. Kan, Z. Xu and C. Gao, Macromolecules, 2011, 44, 444. 94. P. Li, Y. Chen, J. Zhu, M. Feng, X. Zhuang, Y. Lin and H. Zhan, Chem. -­Eur. J., 2011, 17, 780. 95. Y. Deng, Y. Li, J. Dai, M. Lang and X. Huang, J. Polym. Sci., Part A: Polym. Chem., 2011, 49, 1582. 96. B. D. Wang, J. Yang, Z. Zhang, C. Xi and J. Hu, J. Phys. Chem. C, 2011, 115, 24636. 97. X. Ma, F. Li, Y. Wang and A. Hu, Chem.–Asian J., 2012, 7, 2547. 98. J. Liu, C. Chen, C. He, J. Zhao, X. Yang and H. Wang, ACS Nano, 2012, 6, 8194. 99. S. Chatterjee, R. K. Layek and A. K. Nandi, Carbon, 2013, 52, 509. 100. M. Fang, K. Wang, H. Lu, Y. Yang and S. Nutt, J. Mater. Chem., 2009, 19, 7098. 101. R. K. Layek, S. Samanta, D. P. Chatterjee and A. K. Nandi, Polymer, 2010, 51, 5846. 102. M. Fang, K. Wang, H. Lu, Y. Yang and S. Nutt, J. Mater. Chem., 2010, 20, 1982. 103. S. H. Lee, D. R. Dreyer, J. An, A. Velamakanni, R. D. Piner, S. Park, Y. Zhu, S. O. Kim, C. W. Bielawski and R. S. Ruoff, Macromol. Rapid Commun., 2010, 31, 281. 104. L. Ren, S. Huang, C. Zhang, R. Wang, W. W. Tjiu and T. Liu, J. Nanopart. Res., 2012, 14, 940. 105. G. Goncalves, P. A. A. P. Marques, A. Barros-­Timmons, I. Bdkin, M. K. Singh, N. Emamic and J. Gracioa, J. Mater. Chem., 2010, 20, 9927. 106. R. Brako, D. Sokcevic, P. Lazic and N. Atodiresei, New J. Phys., 2010, 12, 113016. 107. J. K. Lee, K. B. Smith, C. M. Hayner and H. H. Kung, Chem. Commun., 2010, 46, 2025. 108. F.-­Y. Sua, Y.-­B. He, B. Li, X.-­C. Chen, C.-­H. You, W. Wei, W. Lv, Q.-­H. Yang and F. Kang, Nano Energy, 2012, 1, 429. 109. R. Mo, D. Rooney, K. Sun and Y. Yang, Nat. Commun., 2017, 8, 13949. 110. N. Maity, R. Ghosh and A. K. Nandi, Langmuir, 2018, 34, 7585–7597.

Chapter 2

Covalent Functionalization of Polymers 2.1  Covalent Functionalization Here graphene oxide (GO) is taken as the precursor and it is produced by chemical oxidation of graphite with a KMnO4/NaNO3 mixture in a concentrated H2SO4 medium using Hummer's method.1 The GO sheets are collected by ultra-­sonication and bulk quantities of GO is thus obtained. rGO is produced by chemical reduction using NaBH4/hydrazine hydrate2 or by thermal reduction (at 950 °C in inert atmosphere for 10 s).3 In GO a sufficient amount of carboxyl, hydroxyl and epoxy groups are present making it hydrophilic in nature, so they are immiscible with most polymers, which are usually hydrophobic in nature. The functional groups present in GO prompt an opening to chemically functionalize with polymers; however, rGO has the deficiency of sufficient functionality for chemical functionalization. Therefore, covalent attachment of small molecule having functional groups create sufficient functionality in GO and rGO to provide outstanding prospects for modification with polymer by a grafting procedure. The grafting of polymer chain either from GO/rGO or from small molecular functionalized GO/rGO can be performed using the ‘grafting to’ or by ‘grafting from’ procedures.

2.2  ‘Grafting to’ Method In this grafting to method the polymer chains are first synthesized and these pre-­synthesized polymers are then attached to functional groups of GO or rGO or to its aromatic surface. The esterification, amidation, nitrene   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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Covalent Functionalization of Polymers

25

chemistry, click chemistry, radical addition, etc. techniques are used in the ‘grafting to’ procedure to covalently link the functional polymers on the GO surface.

2.2.1  Esterification Reaction Poly(vinyl alcohol) (PVA) was functionalized from the GO surface by direct esterification reaction between the carboxylic acid group of GO and the hydroxyl groups of PVA in presence of N,N′-­dicyclohexylcarbodiimide (DCC) and 4-­dimethylaminopyridine (DMAP) catalyst by Salavagione et al.4 The PVA attached GO was then reduced using hydrazine hydrate to produce rGO hybrid (Figure 2.1) which is soluble in water and in DMSO in warm conditions. Here, the degree of esterification is dependent on the PVA tacticity which controls the crystallinity, glass transition temperature (Tg) and thermal stability of the polymer. The segmental motion of PVA chains in the rGO– PVA conjugate having 10% (w/w) PVA content becomes averted due to the

Figure 2.1  Schematic  illustration of the esterification of graphite oxide with PVA and its reduction with hydrazine hydrate (DCC = N,N′-­dicyclohexylcarbodiimide DMAP = 4-­dimethylaminopyridine). Reproduced from ref. 4 with permission from American Chemical Society, Copyright 2009.

26

Chapter 2

presence of rigid graphite structure causing an increase in Tg by 35 °C. The crystallinity of PVA (51% crystallinity) is reduced significantly and the amorphous nature arises due to steric hindrance of GO. However, the degradation temperature of PVA increased by 100 °C as GO acts as a shield for heat flow. In another approach they treated GO with SOCl2 converting it into acid chloride and coupled it with hydroxyl group of PVA. This PVA grafted GO on reduction with hydrazine hydrate exhibited similar changes in properties as delineated above.4 GO is functionalized with poly(vinyl chloride) (PVC) by the same group5 by introducing a hydroxyl group on the PVC backbone by nucleophilic substitution and followed by esterification reaction. This PVC functionalized graphene also exhibits an increase in Tg by 30 °C for the 1.2 wt% GO content, representing that GO hinders the polymer chain mobility. This hybrid exhibits a noteworthy reinforcement property causing an increase in the storage modulus by 70%. The functionalization of GO with PVA by esterification reaction is also made by another group6 producing a mechanically strong elastic thin film of GO, which also exhibits an intense shift in Tg (70 °C to 90 °C) suggesting that attachment of GO nanosheets on the PVA backbone reduces the polymer chain mobility significantly. The grating of water-­soluble polysaccharides, e.g. chitosan (LMC) and hydroxypropyl cellulose (HPC) is made by grafting onto the GO surface by esterification. It shows an attachment of 20% LMC and 30% HPC respectively from the TGA analysis.7 Raman spectra of the sample indicate that the intensity ratio of the D-­and G-­bands (ID/IG) has increased to some extent, suggesting the defect of GO becomes enhanced due to the covalent functionalization. The hydroxyl terminated poly(3-­hexyl thiophene) is grafted with a carboxylic acid group of GO by esterification reaction, producing a donor (polythiophene)–acceptor (graphene) nanohybrid.8 The ensuing P3HT-­grafted GO sheets have good solubility in organic solvents (e.g., THF), enabling easy photovoltaic device fabrication with fullerene. The resultant device exhibits a 200% increase in the power conversion efficiency with respect to the P3HT/C60 counterpart because GO also acts as an acceptor in addition to C60. The esterification procedure is also used to covalently functionalize GO with poly(piperazine spirocyclic pentaerythritol bisphosphonate) (PPSPB). The reduced form between GO-­PPSPB is used to prepare nanocomposites with ethylene vinyl acetate copolymer (EVA) showing enhanced flame retardancy and improved thermal stability9 Turlakov et al.10 used the esterification reaction of GO with the four phenyleneethynylene (PPEs) copolymers by applying the DCC protocol with the addition of DPTS/4-­Py in DMF to promote a better activation of the highly steric carboxylic acids (Figure 2.2). DMF is a good solvent for dispersing GO, but not for PPEs, so to avoid PPEs aggregation during the GO-­PPE formation, they sonicated DMF dispersion, and then the other reactants are added in dichlorobenzene. The dark solutions showed fluorescence (by exciting at 365 nm), which completely disappeared at the end of the reaction. In this system functionalization mainly occurred at the edges of the GO sheets, hence GO-­sheets are not totally

Covalent Functionalization of Polymers

27

Figure 2.2  Reagents  and conditions: DCC/DPTS/4-­PPy, DMF/DCB, 300 W, 4.75 GHz,

600 rpm, 90 min, 150 °C. Reproduced from ref. 10 with permission from Elsevier, Copyright 2017.

covered as the trend of PPEs is to self-­assemble in an ‘edge-­on’ conformation, i.e. the conjugated backbones are parallel to the GO surface, rather than ‘face-­on’, where the backbones are lying flat on the GO surface. This strange supramolecular organization can affect the energy transfer in solid state, hence its device performance. Feng et al.11 used the direct esterification reaction between graphene oxide (GO) and hydroxylated sulfonated poly(ether ether ketone) (SPEEK-­OH) by dispersing them in dimethyl formamide followed by mixing and sonication vigorously for three days with N,N-­dicyclohexylcarbodiimide (DCC), 4-­dimethylaminopyrydine (DMAP) to get the functionalized product of GO (SGO) (Figure 2.3). This is then mixed with sulfonated polyarylene ether nitrile (SPEN) to explore the effect of loading on the performance of proton exchange membranes. It is observed that SGO is uniformly dispersed producing a perpendicular packing structure in the matrix. This causes the formation of long and continuous proton transfer channels effectively improving the proton conductivity.

28

Chapter 2

Figure 2.3  Schematic  diagrams of the synthesis of SPEEK–OH grafted GO. Reproduced from ref. 11 with permission from Elsevier, Copyright 2018.

Figure 2.4  The  schematic illustration of (a) functionalization including the ester-

ification reaction between GO and epoxy and the 3D structure of EGO-­ CNT formed by π–π stacking interaction among EGO and CNTs and (b) the crosslinking network of EGO-­CNT/EP. Reproduced from ref. 12 with permission from the Royal Society of Chemistry.

Qi et al.12 grafted epoxy resin on GO via esterification reaction by dispersing GO in DMF by sonication. An epoxy resin, WSR618, with an epoxy equivalent of 185–192 g eq.−1 was also dissolved in DMF followed by mixing of the two DMF dispersions and was stirred for 1 h at room temperature. Then triphenyl phosphine, TPP, (0.5 wt% to WSR 618) is added as catalyst and the reaction is allowed to proceed for 24 h at 100 °C (Figure 2.4).

Covalent Functionalization of Polymers

29

The reaction product is diluted and filtered with a membrane of 10 mm pore size. Then the obtained solid is washed with excess ethanol five times and dried in vacuum to get epoxy functionalized GO (EGO). By mixing with a 1D carbon nanotube (CNT) a 3D hybrid structure is produced through π-­stacking interaction. The results show that functionalized GO aids good dispersion of CNTs in an epoxy matrix, promoting excellent mechanical properties in the epoxy composites filled with a 3D nanostructure. In the 3D filler network, crack deflection/bifurcation induced by functionalized GO makes a twisted crack path, which makes the cracks encounter more CNTs, resulting in more energy dissipation. This is the probable mechanism for its excellent reinforcing effect. Zhang et al.13 used an environmentally friendly route to synthesize lipophilic GO derivative by esterification reaction of its carboxylate salt and 1-­bromohexadecane (HD) in the presence of a phase-­transfer reagent in water. The long alkyl chains grafted GO sheets (GO-­HD) show the long-­term stable dispersion capability in many organic solvents. The GO-­HD acts as a reinforcing filler to produce thermoplastic polyurethane (TPU) nanocomposites through solution casting method. For the good compatibility of long alkyl chains of GO-­ HD with TPU chains uniform distribution of the filler is noticed due to good interfacial adhesion. Tensile tests indicates that GO-­HD/TPU nanocomposites at low filler content (2 wt%) exhibit higher modulus, ultimate tensile strength from those of GO/TPU nanocomposites made at the same filler concentration. The improved load transfer efficiency is due to load transfer process at a microscopic level, as evident from Raman spectroscopy analysis.

2.2.2  Amidation Reaction Liu et al. produced polyethylene glycol (PEG) grafted nanographene oxide (NGO) from tiny portions of GO by carbodiimide catalysed amidation reaction between NGO and PEG-­amine stars (Figure 2.5). The product PEGylated NGO exhibits outstanding stability in solutions and has been used for the delivery of an insoluble camptothecin (CPT) analogue anticancer drug (SN-­38) which binds by π-­stacking interactions.14 They have found that SN-­38 on NGO-­PEG showed negligible release from NGO in PBS and ∼30% release in mouse serum in three days. This suggests strong noncovalent binding of SN-­38 on NGO sheets and the slow but finite release in mouse serum is caused by binding of SN-­38 by serum proteins.15 Jin et al. also synthesized different molecular weight PEGylated GO nanosheets by amide formation of GO sheets with PEGdiamin. The PEGylated GO nanosheets with free amines could selectively improve trypsin activity and thermostability on three important serine proteases (chymotrypsin, trypsin and proteinase K) while exhibiting no effect on proteinase K or chymotrypsin.16 The authors claimed that this is the first success of functionalized GO as an efficient enzyme positive modulator with high selectivity, exhibiting a novel potential of GO, when appropriately functionalized, in enzyme engineering as well as in enzyme-­based biosensing and detection.

30

Chapter 2

Figure 2.5  Schematic  presentation of SN38 loaded NGO-­PEG. Reproduced from

ref. 14 with permission from American Chemical Society, Copyright 2008.

Cai et al.17 used amidation reaction between a carboxylic acid group of GO and an amine group of poly(ethylene amine) (PEI) to produce water soluble PEI-­GO using carbodiimide catalyst (Figure 2.6). This PEI-­GO has been used to produce graphene/Ag nanocomposite to increase the stability and diminish the cytotoxicity of Ag nanoparticles. This composite shows higher antibacterial activity from that of poly(vinyl pyrrolidone) (PVP)-­stabilized AgNP, and the AgNPs on PEI-­GO are more stable than that of the AgNPs on PVP, ensuring a long-­term antibacterial effect. Shan et al.18 covalently functionalized GO sheets with biocompatible poly-­l-­lysine through an amide linkage between poly-­l-­lysine and GO in an alkaline solution. A schematic representation of the formation of PLL-­functionalized graphene is shown in Figure 2.7 where the PLL is covalently grafted to graphene through amide bond formation by the reaction of epoxy groups of GO and amino groups on PLL in the presence of KOH. This GO-­PLL assembly is further conjugated with horseradish peroxidase (HRP) for the construction of chemically modified gold electrode for H2O2 sensing. Wang et al.19 have covalently grafted epoxy-­based precursor polymer (BPAN) with GO through an ester linkage and have functionalized it with

Covalent Functionalization of Polymers

31

Figure 2.6  The  synthesis of a water-­soluble PEI-­rGO–AgNP hybrid: step 1, oxida-

tive treatment of graphite yields single-­layer GO; step 2, the amidation reaction between carboxylic groups of GO and amine group of PEI to synthesize PEI-­GO; step 3, the ultrasonic treatment of PEI-­GO in the presence of silver nitrate produces a PEI-­GO/Ag+ mixture; step 4, the chemical reduction of GO and silver nitrate with hydrazine monohydrate produces a water-­soluble PEI-­rGO–AgNP dispersion. Reproduced from ref. 17 with permission from Elsevier, Copyright 2012.

Figure 2.7  Schematic  diagram of graphene-­PLL synthesis (green line represents

PLL). Reproduced from ref. 18 with permission from American Chemical Society, Copyright 2009.

hyperbranched azo-­polymer by azo-­coupling reaction with hyperbranched diazonium salts under extremely mild conditions. From the XPS study the degree of functionalization is measured to be approximately one azobenzene repeat unit per 30 carbons of GO. Qin et al.20 covalently functionalized GO with 3-­aminopropyl pyrrole (APP) monomer by reacting the carboxyl groups of GO and the amine groups of

Chapter 2

32

APP. Then, a GO/poly(1-­(3-­aminopropyl)-­pyrrole) (poly(GO–APP)) copolymer was synthesized by in situ crosslinking polymerization using ammonium persulphate as the initiator (Figure 2.8). Here GO acts as a crosslinking agent like a polyfunctional monomer. This poly(GO–APP) copolymer has been used for the adsorption of Cu2+, Ni2+, Pb2+ and Cd2+ metal ions and its adsorption capacity has highly increased compared to GO. Amongst the four metal ions Cu2+ is most effectively adsorbed by poly(GO-­APP). Besides the greater adsorption capacity for metal ions, the poly(GO–APP) shows a good conducting property where pure GO is an insulator. This is due to covalent functionalization with the conjugated π–π system present in poly(GO–APP) expanding its uses in different fields. Wu et al.21 chose GO as graphene precursor for the production of hyperbranched aromatic polyamide functionalized graphene sheets (GS–HBA) by amidation reaction (Figure 2.9).

Figure 2.8  Synthesis  process of the poly(GO–APP) copolymer. The inset photos show the GO NMP solution, GO–APP NMP solution, poly(GO–APP) NMP solution and gelatinous poly(GO–APP) after filtration (from left to right). Reproduced from ref. 20 with permission from the Royal Society of Chemistry.

Figure 2.9  Synthesis  of HBA functionalized graphene sheets. Reproduced from ref. 21 with permission from the Royal Society of Chemistry.

Covalent Functionalization of Polymers

33

GO is first modified by surface linkage of ethylene diamine(EDA) via amidation reaction, rendering GO to have enough active sites for initiating the hyperbranching reaction. The modified GO is then grafted with hyperbranched aromatic polyamide by using 3,5-­diaminobenzoic acid (DABA). Finally, the resulting HBA-­functionalized GO is chemically reduced by means of high temperature in N-­methyl pyrrolidone (NMP) to obtain GS–HBA. Hyperbranched polymer functionalized GO significantly improves the dispersion of GS in the thermoplastic polyurethane matrix and reinforces the interfacial adhesion between GS and TPU matrix at molecular level resulting in a higher efficient load transfer between the polymer phase and GSs. Further the mechanical properties of GS–HBA composites can be tuned from an elastomer to a rigid and ductile thermoplastic by successive addition of GS–HBA. The HBA chains can not only improve dispersion and compatibility of GS in the TPU matrix, but also serve as dielectric layers; as a result, the GS–HBA–TPU composites exhibit superior dielectric performance and higher permittivity.21 Xu et al.22 covalently grafted six-­armed poly(ethylene glycol) on GO sheets via a facile amidation reaction under mild conditions, making GO-­PEG (PEG: 65 wt%, size: 50–200 nm), stable and biocompatible in physiological solution (Figure 2.10). This nanosized GO-­PEG is nontoxic to human lung cancer A549 and human breast cancer MCF-­7 cells. The paclitaxel (PTX), a widely used cancer chemotherapy drug, is coupled onto GO-­PEG surface via π–π stacking with a high loading capacity of PTX (11.2 wt%). This complex can quickly enter into A549 and MCF-­7 cells, as evident from fluorescence microscopy, and it shows very high cytotoxicity to MCF-­7 and A549 cells in a wide range of PTX concentration and time from that of free

Figure 2.10  Structure  of six-­armed poly(ethylene glycol) grafted on GO. Reproduced from ref. 22 with permission from American Chemical Society, Copyright 2014.

Chapter 2

34

PTX. This nanoscale drug delivery system on the basis of PEGylated GO may find potential application in biomedicine. Li et al.23 have prepared primary antibody-­AFP (Ab1) functionalized reduced rGO via an amidation reaction between the carboxylic acid group of the rGO and amine groups of Ab1. They have developed a nonenzymatic electrochemical immunosensor based on palladium nanoparticles conjugated on the above functionalized rGO (Pd-­GS) for sensitive detection of a cancer biomarker by one-­spot synthesis to immobilize secondary antibody (Ab2). The resulting Pd-­GS-­Ab2 conjugate acts as immunosensor to detect AFP and it shows good recovery in the assay results for AFP in human serum samples.

2.2.3  Click Chemistry In another method of the ‘grafting to’ approach click chemistry technique is used to graft azido terminated polymer chains with alkyne derivative of graphene. Pan et al.24 grafted poly(N isopropylacrylamide) (PNIPAm) to produce water-­soluble graphene for drug delivery by this method. On this occasion, they modified PNIPAm with an azide group by polymerizing the monomer NIPAm by ATRP reaction and substituting its halide end-­groups with azide groups, as shown in Figure 2.11.

Figure 2.11  Synthetic  scheme of PNIPAM-­GS using click chemistry. (ECP = ethyl

2-­chloropropionate; IPA = isopropanol; Me6tren = Tris[2-­(dimethlyamino) ethyl]amine; PMDETA = N,N,N′,N″-­pentamethyl diethylenetriamine). Reproduced from ref. 24 with permission from John Wiley & Sons, Copyright 2011 WILEY‐VCH Verlag GmbH & Co. KGaA, Weinheim.

Covalent Functionalization of Polymers

35

Then they synthesized an alkyne derivative of GO using an amide linkage and coupled the two following the scheme presented in the Figure 2.11. The pure PNIPAm homopolymer has a lower critical solution temperature (LCST) 37.8 °C but in the PNIPAm functionalized GO it shows LCST at 33 °C, which is ∼5 °C lower than that of pure PNIPAM. This lowering of LCST is attributed to the hydrophobic nature of GO attached to PNIPAM making it have lower solubility in water. The GO functionalized with PNIPAm can load a hydrophobic anticancer drug (camptothecin, CPT) with a superior loading capacity due to hydrophobic and π–π stacking interaction between PNIPAm functionalized GO and the aromatic drug. The in vitro drug release experiment has exhibited 16.9% and 19.4% release of CPT in water and PBS buffer (pH7.4) after 72 h at 37 °C. The PNIPAm functionalized GO does not possess any practical toxicity, so it acts as an effective delivery vehicle for anticancer drugs. Jin et al.25 covalently grafted poly(ethylene glycol) (PEG) from the graphene surface obtained by the CVD method by dispersing in water using sodium dodecylsulfate (SDS), and then adding 4-­propargyloxybenzenediazonium tetrafluoroborate into the mixture followed by reaction at 45 °C for 8 h. This produces alkynyl-­functionalized graphene flakes which was coupled with polyethylene glycol with azido and carboxyl end-­groups (azido-­dPEG4-­acid) (Figure 2.12). The ID/IG value increased from ∼0.10 to ∼0.71,which indicates that the reactivity of CVD-­graphene is higher than that of the isolated graphene flake. Sun et al.26 also synthesized azide terminated polystryrene (PS), firstly prepared by ATRP initiated by 3-­azidoethyl 2-­bromoisobutyrate and then GO sheets are functionalized with propargyl alcohol by an acylation reaction, followed by click reaction at room temperature with CuBr/PMDETA as the catalyst in DMF between alkyne functionalized graphene sheets and azido-­ terminated PS. The azide terminated polystyrene, poly(methyl methacrylate) (PMMA), poly (4-­vinyl pyridine) (P4VP), poly(methyl acrylic acid) (PMAA),

Figure 2.12  Diazonium  reaction and subsequent click chemistry functionalization

on graphene sheets. Reproduced from ref. 25 with permission from American Chemical Society, Copyright 2011.

36

Chapter 2

poly(dimethylamino ethyl methacrylate) (PDMA) polymers were synthesized using the RAFT polymerization technique with azide-­terminated chain transfer agents and they were subsequently grafted onto alkyne-­functionalized graphene using click chemistry.27,28 Kwon et al.29 recently prepared poly(sodium 4-­styrenesulfonate) (PSS) using RAFT polymerization. The benzodithioate part in this polymer is transformed to an end functional thiol group using a reducing agent. The RAFT polymer is then attached to GO using a thiol–ene click reaction. TEM and AFM results show that the produced functionalized GO (f-­GO) does not re-­ aggregate in water, even when it was reduced to functionalized reduced GO (f-­rGO). Cernat et al.30 have inserted the azide group into the GO backbone by chemical functionalization on mixing GO and sodium azide dispersions in water and is kept under continuous stirring at 10 °C for 1 h. After freeze-­ drying the reaction mixture the product is dispersed in 100 mL water at 10 °C and then purified by centrifuging. The final product was isolated by freeze-­ drying and characterized by spectroscopy and microscopic techniques. After the successful synthesis of the graphene-­azide platform it is reacted by clicking ethynyl ferrocene electrochemically. Wang et al. reported a methodology that utilizes graphene derivative functionalized with azides as the click reagent and it quickly couples to alkyne-­modified DNA strands to create stable polyvalent conjugates with exceptionally high DNA densities on graphene nanosheets (Figure 2.13). In this approach DNA–graphene conjugates exhibit excellent integration and stability of assembly with other multiple DNA nanostructures making advanced DNA nano-­architectures on the 2D platform of graphene. Punetha et al.,32 in a recent review, recorded the use of click chemistry in the functionalization of graphene and carbon nanotubes not only for polymers but also with biomolecules and nanoparticles for different biotechnological and energy applications. In a word all types of molecules can be grafted in this click chemistry grafting approach. The solubility of the polymer functionalized graphene can be adjusted from water-­soluble to oil-­soluble, acidic to basic and polar to apolar by selecting a suitable polymer for grafting. This click chemistry approach is easily feasible with mild reaction conditions and has good control of the grafted polymer structure.

2.2.4  Nitrene Chemistry The nitrenes are highly reactive intermediates and are produced from azide groups through either thermolysis or irradiation, which has been tried for the covalent functionalization of CNTs, fullerene, etc. via the [2+1] cycloaddition of nitrenes to the π-­electron system.33–35 Nitrene addition as a catalyst is used in click chemistry when it is hard to remove copper from the hybrid. He et al.36 used this technique to graft azido terminated polystyrene (PS) and to graft azido terminated polyethylene glycol (PEG) chains on the graphene surface (Figure 2.14).36

Covalent Functionalization of Polymers

37

Figure 2.13  Schematic  representation of the synthesis route of DNA–graphene conjugates Reproduced from ref. 31 with permission from the Royal Society of Chemistry.

PS grafted graphene shows good dispersibility in DMF, THF, toluene, chloroform, etc. while PEG grafted graphene exhibits full exfoliation and good dispersibility in water. This strategy permits different kinds of polymers and functional moieties to be covalently bonded on graphene, producing functionalized graphene nanosheets. The functional groups introduced on graphene promote much better thermal and mechanical properties than those of GO, and can be further modified by differently chemical reactions, including amidation, surface-­initiated polymerization and reduction of metal ions. Xu et al.37 produced conjugated polyacetylene functionalized graphene utilizing the nitrene chemistry through the reaction between the reactive azide groups and graphene moieties (Figure 2.15) and it shows good dispersibility in various organic solvents.37 Similar emission curves, with an identical emission maximum suggest that the attachment of polyacetylene chains to graphene almost did not affect their electronic structure; however, their fluorescent quantum yields are somewhat higher than that of polyacetylene.

38

Chapter 2

Figure 2.14  General  strategy for the preparation of functionalized graphene

nanosheets (f-­GNs) by nitrene chemistry and the further chemical modifications. Reagents and conditions: (i) NMP, 160 °C, 18 h; (ii) ε-­caprolactone, stannous octoate, 120 °C, 24 h; (iii) palmitoyl chloride, TEA, r.t., 24 h; (iv) styrene, CuBr/PMDETA, 80 °C, 24 h; (v) FeCl3, NaOH, DEG, 220 °C, 1 h. Reproduced from ref. 36 with permission from American Chemical Society, Copyright 2010.

Figure 2.15  Synthetic  root of polyacetelene–graphene composites using nitrene

chemistry [Pac = Poly{[1-­(4-­tolyl)-­5-­chloro-­1-­pentyne]-­co-­[1-­(4-­tolyl)-­ 5-­azido-­1-­pentyne]}]. Reproduced from ref. 37 with permission from John Wiley & Son, Copyright 2011 WILEY‐VCH Verlag GmbH & Co. KGaA, Weinheim.

Xu et al.38 have synthesised cyclodextrin functionalized rGO via simultaneous reduction of GO to rGO and grafting of mono-­(6-­azido-­6-­deoxy)-­b-­ cyclodextrin (CD) via nitrene addition onto rGO in one-­pot synthesis. This is because the solvothermal reduction of GO and nitrene addition have the same reaction condition (Figure 2.16).

Covalent Functionalization of Polymers

39

Figure 2.16  Synthesis  route for the preparation of RGO-­g-­CD/Fc-­PNIPAM and RGO-­ g-­CD/Fc-­HPG nanohybrids. Reproduced from ref. 38 with permission from Elsevier, Copyright 2013.

The lyophilized rGO-­g-­CD powders can be readily dispersed in water, NMP and N,N-­dimethylformamide (DMF) upon ultrasonic agitation. The obtained dispersion is homogeneous and stable. The rGO-­g-­CD nanosheets are capable of further functionalization with ferrocene-­modified hyperbranched polyglycerol (HPG-­Fc) and poly(N-­isopropylacrylamide-­co-­vinylferrocene(ii)) (PNIPAM-­Fc) via host-­guest inclusion complexation of CD and ferrocene (Fc) moieties. Both rGO-­g-­CD/Fc-­PNIPAM and rGO-­g-­CD/Fc-­HPG nanohybrids are dispersible in a wider range of solvents and are stable from that of the rGO-­ g-­CD precursor nanosheets. For the thermoresponsive behaviour of PNIPAM moieties, the rGO-­g-­CD/Fc-­PNIPAM shows a lower critical solution temperature (LCST) of 28 °C giving an opportunity to develop thermoresponsive nanodevices. The rGO-­g-­CD/Fc-­HPG nanohybrids shows low cytotoxicity towards 3T3 fibroblasts in 3-­(4,5-­dimethylthiazol-­2-­yl)-­2,5-­diphenyl tetrazolium bromide (MTT) cell viability assay. Soleimani et al.39 synthesized a GO-­cellulose nanowhisker nanocomposite hydrogel using nitrene chemistry. The obtained porous nanocomposite hydrogel can efficiently remove cationic dyes such as methylene blue (MB) and Rhodamine B (RhB) from wastewater with high absorption power as evident from UV–vis spectra. The adsorption process have shown that 100% of MB and 90% of RhB are removed and the equilibrium state is reached in 15 min for low concentration solutions in accordance with the pseudo-­second-­ order model. Moreover, the sample shows recyclability after being used several times. High adsorption capacity and easy recovery make these graphene grafted materials good adsorbent materials for water pollutants and wastewater treatment. Hu et al.40 synthesized a series of crosslinking composite fuel cell membranes with quaternized polysulfone (QPSU) and rGO using nitrene chemistry. A crosslinking structure between QPSU chains containing azide groups

Chapter 2

40

Figure 2.17  Synthetic  process of azide-­assisted crosslinked composite membranes

Reproduced from ref. 40 with permission from Elsevier, Copyright 2017.

and rGOs is formed by thermal activation without using extra crosslinking agents (Figure 2.17). Nitrene addition avoids the copper catalyst used in click chemistry where it is hard to remove copper which affects the performance of membranes. The covalent crosslinking fixes the rGO well and improves the compatibility of rGO and QPSU. The stable microstructure of the membrane after crosslinking affords good chemical stability of membranes and relatively high conductivity, and regulates the fuel permeation.

2.2.5  Radical Addition GO sheets are nicely exfoliated in a reactive poly(methyl methacrylate) (PMMA) matrix produced by in situ nitroxide mediated polymerization (NMP) producing end-­capped by a cleavable alkoxyamine (PMMA-­ONR2). Heating the mixture at 50 °C under extensive stirring leads to the C–O bond cleavage of the alkoxyamine producing PMMA macroradicals that add to sp2 carbons of GO. Thus in situ reduction of GO with the simultaneous grafting of PMMA is made by a radical addition pathway under mild biphasic conditions, producing graphene grafted PMMA.41 This grafting, which occurs during GO reduction, avoids the re-­stacking of rGO and leads to highly conductive rGO– PMMA nanocomposites under mild experimental conditions. Shen et al.42 in situ polymerized acrylic acid (PAA) and polymerized acryl amide (PAM) in the presence of GO using (NH4)2S2O8 as initiator and graphene grafted polymer is formed by a radical coupling reaction. The polymer modified GO exhibits high quality dispersion in water and the PAM and PAA modified graphene introduce positive and negative charges respectively on the graphene surface. They self-­assemble by a layer-­by-­layer technique via electrostatic interaction to form a multilayer structure.

Covalent Functionalization of Polymers

41

Figure 2.18  Procedures  used in the preparation of graphene/PS–PAM (BPO = ben-

zoyl peroxide). Reproduced from ref. 43 with permission from John Wiley & Sons, Copyright 2009 Wiley‐VCH Verlag GmbH & Co. KGaA, Weinheim.

Shen et al.43 synthesized GO grafted copolymer of styrene and acryl amide (GO-­g-­PS–PAM) by reacting GO with styrene and acryl amide using benzyl peroxide initiator (Figure 2.18). Here the propagating PS–PAM copolymer macroradicals add to the layer of graphene sheets until the copolymer molecules completely wrap around the surface of graphene nanoplatelets. This modified graphene is dispersible with a wide variety of solvents and polymer matrices and could be used as starting material for the fabrication of graphene composites. Recently, Kan et al.44 polymerized different monomers on a GO surface using azobisisobutyro nitrile (AIBN) as initiator where the chain radical combines with graphene. After initiating vinyl monomers by free radical polymerization (FRP), where macromolecular radicals are formed immediately and part of the macromolecular radicals adds to double bonds of GO, producing graphene oxide-­based polymer brushes and simultaneously generates new radicals at its surface, responsible for further radical propagation, termination or transfer (Figure 2.19). The GO-­ polymer bushes do not self-­aggregate; e.g. GO-­g–PGMA shows a very low intrinsic viscosity in DMF solution (∼100 mL g−1), much lower than that of pristine GO dispersion (∼780 mL g−1) and this is close to the value of a common linear polymer. This result indicates that the GO-­polymer brushes do not possess intermolecular chain entanglements and are similar to the globular macromolecules enabling the 2D brushes to be useful as nanofillers to improve the processability and performance of common polymers. In a similar way, stimuli-­responsive polymers, e.g. poly(acrylic acid) and poly (N-­isopropylacrylamide), are covalently grafted onto the GO surface through a facile redox polymerization initiated by cerium ammonium nitrate in aqueous solution at mild temperature.45 The GO-­PAA and GO-­PNIPAm in aqueous solutions could assemble and disassemble by varying the pH and temperature of the solutions respectively. Reduction of GO with dopamine can produce self-­polymerization of dopamine to produce polydopamine (PDA) causing simultaneous capping of GO by polydopamine, increasing the thermal stability.46 The PDA-­capped rGO

42

Chapter 2

Figure 2.19  Synthesis  of 2D macromolecular brushes by free radical polymeriza-

tion of various monomers with the backbone of GO sheets (the blue sheet represents GO, and the red dots and double bonds represent radicals and active grafting points, respectively.) Reproduced from ref. 44 with permission from American Chemical Society, Copyright 2011.

reacts with thiol-­terminated or amino-­terminated poly(ethylene glycol) (PEG) and produces PEG functionalized PDA-­rGO using the ‘grafting-­to’ process. This PEG-­functionalized rGO exhibits good quality of dispersion both in organo-­ and aqueous medium.46 Deng et al.47 synthesized PNIPAM-­ functionalized graphene sheets by pre-­synthesizing PNIPAM using atom transfer nitroxide radical coupling (ATNRC) followed by polymerization onto the graphene surface. It exhibits excellent dispersibility in various organic solvents and in water it shows a temperature responsive behaviour (LCST) ∼37 °C for the presence of thermoresponsive PNIPAM chains. Zhang et al.48 synthesized GO-­grafted poly(vinyl acetate) (PVAc) via a facile approach by γ-­ray irradiation-­induced graft polymerization. The GO-­g-­PVAc forms extremely stable dispersion in common organic solvents and exhibits a great potential in producing graphene-­based composites by solution-­processes. Similarly, Chen et al.49 made GO-­g-­polystyrene by successive intercalation, grafting and exfoliation of graphite oxide in styrene monomer followed by γ-­ray irradiation. Roppolo et al.50 used light induced free radical polymerization on a GO surface by first grafting benzophenone (BP), which later acts as initiator for UV-­grafting of different monomers on the GO surface. This procedure also allows the simultaneous UV-­reduction of unreacted functional groups of GO during the first grafting step where UV excited BP generates semipinacol radicals in solution that recombine with the residual radicals on the GO sheets creating the covalent bond. At the same time, GO undergoes a light induced

Covalent Functionalization of Polymers

43

reduction generating rGO. This procedure leads to the implementation of a methodology for graphene functionalization, with a great variety of monomers (ethyleneglycol methacrylate, EGMA, perfluoro butyl acrylate, PFBA; 2-­(dimethylamino)ethyl methacrylate DMAEM, etc.), exploiting the photosensitive properties of GO-­BP. These grafted materials show improved dispersion into organic solvents and polymeric matrices, which is used for the fabrication of high performing polymer nanocomposites. Skaltsas et al.51 used a facile approach for the covalent functionalization and exfoliation of graphite by one step in situ free radical polymerization of three different monomers, vinyl-­benzyl chloride, styrene, and vinyl-­benzyl trimethylammonium chloride, in the presence of AIBN radical initiator. The polymer-­functionalized graphene sheets are easily soluble in common organic solvents and water, depending on the nature of the grafted polymer. The sonication of graphite affects the quality of exfoliated sheets, as observed from the incorporation of defects and oxygenated species on the graphitic basal plane; the particular one-­pot synthetic method prevents such handicaps. Voylov et al.52 polymerized sodium 4-­vinylbenzenesulfonate in GO dispersion in argon at 75 °C for 18 h. The yield of polymerization of sodium 4-­vinylbenzenesulfonate with GO shows a much greater value (21–90%) than that of the control solution (8.5–11%), i.e. without GO. The polymerization of the control (without GO) occurs by thermally induced autopolymerization. The dependence of yield and molecular weight of the resulting poly(sodium 4-­vinylbenzenesulfonate) (PSSNa) on GO concentration indicates that GO acts as free radical initiator in the thermal polymerization. Kundu et al.53 grafted PNIPAM onto GO dispersion in water via in situ free radical polymerization using ammonium persulphate as initiator under nitrogen atmosphere at 65 °C for 48 h. The product, collected by centrifugation, is repeatedly washed with cold water to remove the noncovalently attached free polymers to GO. The resulting product PNIPAM-­grafted GO (GPNM) exhibits enhanced thermal stability, improved dispersibility both in aqueous and in cell medium, and better biocompatibility and cell viability compared to GO. GPNM exhibits an exciting fluorescence property in aqueous medium with an increase in intensity at 36 °C due to the lower critical solution temperature (LCST) of PNIPAM chains. Moreover both hydrophilic (doxorubicin (DOX)) and hydrophobic (indomethacin (IMC)) drugs loaded on the surface of GPNM hybrid exhibits its efficacy as an efficient carrier for both types of drugs. Redox systems, with Ce(iv) salts in aqueous solutions, are great initiators for vinyl polymerization and homopolymers, block copolymers, and graft copolymers can be easily formed with this system. Suitable reducing agents include alcohols, aldehydes, ketones and amines. Ma et al.54 showed that hydroxyl groups on GO surfaces make perfect reducing agents, if coupled with Ce(iv) salts, for initiation of vinyl monomers, producing polymer brushes grafted on the surface of GO. The redox reaction systems have the prime advantage of operating in aqueous solutions at moderate temperatures hence the process is ‘green’ compared with the polymerization carried

44

Chapter 2

in organic solvents. The initiating sites are present directly on the GO surface, free polymers can be neglected, and the materials are easily separated by filtration or centrifugation. The mechanism of polymer syntheses onto GO sheets by redox polymerization is shown in Figure 2.20a. First, free radicals are produced on GO sheets by the redox reaction of hydroxyl groups

Figure 2.20  (a)  Polymerization mechanism of monomers on the surface of GO

using a Ce(iv) redox initiation system (styrene (St): A = H, B = phenyl; methyl methacrylate (MMA): A = CH3, B = COOCH3; acrylonitrile (AN): A = H, B = CN; acrylamide (AM): A = H, B = COONH2). (b) Structural illustration of two types of polymer grafted GO sheets. Reproduced from ref. 54 with permission from Elsevier, Copyright 2012.

Covalent Functionalization of Polymers

45

of GO sheets with Ce(iv)/HNO3 to initiate the polymerization, then polymer brushes are formed on GO sheets by chain propagation; the termination of the growing chains occurs by coupling with other radicals (mutual termination), disproportionation, oxidative termination with Ce(iv) and chain transfer reactions. Since the initiation of polymerization is from the hydroxyl groups located mainly on the surface of GO sheets, the grafted polymers are expected to be on the surfaces of GO sheets. Two types of structures are expected contingent on whether the layered structures are formed or not, as presented in Figure 2.20b. One is an intercalated structure, where the ordered layer structure remains in the composites after low polymer grafting, and the other is the exfoliated structure, where the ordered layer structure disappears due to the dense polymer grafting, causing random packing of GO sheets. Liu et al.55 reported the liquid crystallinity (LC) of polymer-­grafted GO and its macroscopic assembled nacre-­mimetic composite by grafting polyacrylonitrile (PAN) chains onto GO surfaces via a simple free radical polymerization process. The PAN-­grafted GO (GO-­g-­PAN) sheets dispersed in polar organic solvents, e.g. dimethyl sulfoxide (DMSO) and dimethylformamide (DMF), exhibit nematic and lamellar liquid crystals (LCs) upon increasing concentration. A strong signal present in the circular dichroism spectra of the LCs, suggests the formation of helical lamellar structures of the GO-­g-­PAN LCs. Using industrially viable wet-­spinning technology from the GO-­g-­PAN LCs the obtained fibres exhibit layered structures of GO and PAN, mimicking the classic ‘brick-­and-­mortar’ microstructure present in nacre. The composite displayed excellent mechanical properties with Young's modulus 8.31 GPa, tensile strength 452 MPa and breakage elongation 5.44%. This offers a new approach for the fabrication of ultrastrong and tough biomimic composites. Pooresmaeil and Namazi56 made stimuli-­responsive graphene oxide/polymer brush nanocomposites (GPBNs) prepared through the polymerization of acrylic acid (AA), N-­isopropylacrylamide (NIPAM) and acrylated β-­cyclodextrin (Ac-­β-­CD) from the graphene oxide (GO) surface. Nanocomposites were prepared via free radical polymerization of acrylated β-­cyclodextrin (Ac-­β-­CD) produced from reaction of β-­CD with acrylates. Grafting of Ac-­β-­CD increases the biocompatibility of nanocomposites. With varying temperature or pH of the solution, the assembling and disassembling property of the polymer-­ grafted GO in aqueous solution is controlled, making it useful for controlled drug delivery.

2.2.6  Other Methods The in situ ring-­opening polymerization of caprolactam in the presence of GO, initiated by 6-­aminocaproic acid and the condensation reaction between the active terminal amino groups of nylon chains with the carboxylic acid groups of GO at 250 °C produces nylon functionalized graphene57 (Figure 2.21).

46

Chapter 2

These nylon grafted graphene sheets (NG) fibres made by a melt spinning process, exhibits good compatibility with a nylon matrix producing high performance graphene composite which exhibits increased tensile strength by 2.1-­fold and Young's modulus by 2.4-­fold with a graphene loading of 0.1 wt% only, revealing an excellent reinforcement to composites by graphene. This in situ condensation polymerization technique presents a means to prepare graphene-­based high performance nanocomposites of condensation polymers. Sun et al. synthesized a GO hydrogel by direct cross-­linking GO sheets with poly(N isopropyl acryl amide-­co-­acrylic acid) (PNIPAm-­co-­AA) microgels in water via the reaction between epichlorohydrin (ECH) and carboxyl groups of GO in water.58 This hydrogel shows dual thermal and pH response with good reversibility. Li et al.59 synthesized the ‘charm-­bracelet’-­ type poly (N-­vinyl carbazole) (PVK) functionalized rGO (Figure 2.22) via formation of

Figure 2.21  Synthesis  of NG composites by in situ ring-­opening polymerization of caprolactam in the presence of GO. Reproduced from ref. 57 with permission from American Chemical Society, Copyright 2010.

Figure 2.22  Synthesis  of RGO-­PVK. Reproduced from ref. 59 with permission from John Wiley & Sons, Copyright 2011 WILEY‐VCH Verlag GmbH & Co. KGaA, Weinheim.

Covalent Functionalization of Polymers

47

anions on the PVK backbone in presence of sodium hydride followed by nucleophilic addition into the π-­conjugated ring of the rGO sheets. The solubility of PVK functionalized GO increases drastically and the hybrid exhibits excellent broadband optical limiting responses at 532 and 1064 nm. A series of poly(l-­lactide) (PLLA) thermally reduced graphene oxide (TrGO) composites (GLLA) were synthesized by Yang et al. using in situ ring-­opening polymerization of lactide.60 The conductivity GLLA composites increases with increasing TrGO content and it shows insulating conductive percolation behaviour between 1.00 and 1.50 wt% TrGO content. Poly(bisphenol a-­coepichlorohydrin) (PBE) chains are covalently attached to the GO nanosheets using a carbodiimide activating agent (N,N′-­dicyclohexylcarbodiimide (DCC)) to couple GO carboxylic acid groups and hydroxyl groups of the PBE chains. PBE-­functionalized-­GO (PBE-­GO) has been used to form PBE-­GO : PBE thin-­film nanocomposites by solution mixing of the components followed by solvent evaporation and it exhibits very high levels of mechanical reinforcement.61 Zhao et al.62 grafted 2-­bromopropionyl bromide onto the r-­GO surface utilizing a diazonium addition reaction followed by covalent attachment of poly(1-­ vinylimidazole) (PVI). The as-­s ynthesized PVI grafted rGO was used as an adsorbent to remove methylene blue (MB) from water solution, showing a high adsorption capacity, and thus finds use for the removal of anionic dyes from water solution for sustainability. McGrail et al.63 made rapid covalent functionalization of GO or rGO in acidic aqueous suspensions under ambient conditions utilizing the Pinner reaction between hydroxyl groups on (r)GO and nitriles. The modified platelets have tuneable solubility, as well as multiple functionalities. In this method, GO requires little purification and no drying after preparation and multiple grams can be functionalized with a variety of small molecule and polymer nitriles in only a couple of hours. Maity et al.64 grafted polyaniline by using covalent functionalization of paraphenelyne ethylene diamine to acylated GO which is subsequently used to polymerize aniline monomer in the presence of p-­toluenesulfonic acid. Concomitantly, GO undergoes in situ reduction to a certain extent thereby yielding highly conducting fibrillar networks. Nandi's group also have reported the synthesis and self-­ assembly of graphene quantum dots grafted with poly(ε-­caprolactone) of different degrees of polymerization 3, 7, 15 and 21 produced from ring opening polymerization.65

2.3  ‘Grafting from’ Method In the ‘grafting from’ technique the polymerization of monomers occurs from the macro initiators derived from the surface of graphene. These initiating species are covalently attached directly from the carboxylic acid or hydroxyl groups of GO. Also by initial grafting of the small molecules from GO to convey chosen functionality followed by attachment of the initiator. The major advantages of this ‘grafting from’ technique is that the steric interference

Chapter 2

48 66

cannot limit the polymer chain growth. Liu et al. used graphene peroxide (GPO), produced by γ-­irradiation, for in situ polymerization of acrylamide (AM) by free radical polymerization to produce GO-­g-­PAM. GPO is used as a polyfunctional initiating and cross-­linking centre to produce graphene-­based polymer hydrogels. The hydrogel nanocomposite exhibits medium elastic moduli, very high tensile strengths and extremely high extensibility of 5300%, and also exhibits very low hysteresis and excellent resilience properties. Pristine graphene functionalized conjugated polymers were synthesized by Ma et al.67 using the Bergman cyclization of enediyne-­containing molecules. This conducting polymer grafted graphene samples exhibits good solubility in a variety of organic solvents and also possesses good electrical conductivity. Chatterjee et al.68 grafted amino-­functionalized reduced graphene oxide (a-­ RGO) with polyaniline by polymerizing aniline with ammonium persulphate in acetic acid medium. The morphology of polyaniline (PANI) changes from a nanotube to a flat rectangular nanopipe (FRNP) by the polymerization of aniline and an enormous (∼500 times) improvement in photocurrent is observed in FRNP over PANI nanotubes on irradiation with white light of one sun intensity. The photoresponse is reproducible in cyclic runs with a time interval of 100 s and in both the negative and positive bias the photocurrent increases with increase in bias voltage. The monomer 3-­(2-­hydroxyethy l)-­2,5-­thienylene (HET) is grafted with acyl chloride functionalized reduced graphene oxide (frGO) to form HET-­g-­rGO which on oxidative polymerization with FeCl3 produces poly [3-­(2-­hydroxyethyl)-­2,5-­thienylene] grafted rGO.69 It is used as an efficient substitute active material of TiO2 in the dye sensitized solar cell that yields an overall power conversion efficiency of 3.06% with N-­719 dye. The most important methods employed for ‘grafting from’ techniques are the atom transfer radical polymerization (ATRP) and reversible addition-­fragmentation chain transfer (RAFT), to be discussed below.

2.3.1  Atom Transfer Radical Polymerization (ATRP) Matyjaszewski70 and Sowamoto71 proposed a general scheme for Cu-­based ATRP as follows:

The ATRP method of polymerization consists of a dynamic equilibrium between the alkyl halide (Pn–X) species (dormant chains) and propagating radicals, being established by reversible homolytic halogen transfer between a dormant chain and a transition metal complex at its lower oxidation state. This forms propagating radicals and a higher oxidation state metal halogen complex. At first, an alkyl halide initiator reacts with CuI complex (lower oxidation state) to produce alkyl radical and a CuII based complex (higher oxidation state); the later acts as a persistent radical. Therefore, in

Covalent Functionalization of Polymers

49

the polymerization medium there exists a relatively low concentration of Pn° which reacts with the monomers in the propagation step minimizing terminations. Other transition metal complexes like Fe, Ru, Mn, etc. may also be used; however, Cu-­based ATRP has received maximum attention so far. A large number of reports of polymer functionalization of graphene by the ATRP technique exists in the literature and is flourishing, extensively aimed at different applications. Surface initiated ATRP (SI-­ATRP) can tune the molecular weight and polydispersity index (PDI) of the grafted polymer chain. The GO-­g-­polymer halide group is present at the chain end group, so one can think substitution of halide group by organic functional groups or plan for block copolymer formation using the carbon–halide bond as the initiating site. Thus, ATRP is used for the graphene-­based filler modification and also in situ composite formation when both the graphene grafted polymer and matrix polymer are the same. Fang et al. prepared graphene/polystyrene nanocomposite by in situ ATRP of graphene grafted initiator (Figure 2.23).72 They first introduced the hydroxyl group on the rGO surface by the diazonium coupling reaction between the rGO and 2-­(4-­aminophenyl) ethanol, then covalent attachment of 2-­bromo isobutryl bromide (BIB) is made from the hydroxyl group by forming ester linkage. Polystyrene (PS) chains are then grafted using ATRP technique from the graphene based macro initiator. The glass transition temperature (Tg) of

Figure 2.23  Synthesis  root of polystyrene functionalized graphene nanosheets

(PMDETA = N,N,N′,N″-­pentamethyldiethylenetriamine; MBP = 2-­bromo-­ propionate) Reproduced from ref. 72 with permission from the Royal Society of Chemistry.

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Chapter 2

PS in the grafted state is 15 °C more than that for pure polystyrene. Composites of this functionalized graphene sheet with PS increases the mechanical properties of the PS polymer matrix, e.g. Young's modulus increases 57.2% and tensile strength increases 69.5%, by adding only 0.9 wt% functionalized graphene sheet compared to those of pure PS film. The same group has varied the chain length and graft density of the grafted chains by controlling the concentrations of diazonium compound on the rGO surface and monomer before the ATRP.73 For low molecular weight of grafted PS and high grafting density Tg increases to 18 °C but for high molecular weight of grafted PS and for low grafting density, a smaller increase of Tg (9 °C) occurs, probably due to superior heat conductivity of graphene. Ren et al.74 polymerized PNIPAm in a similar way from the graphene surface and the functionalized graphene is highly temperature sensitive in aqueous medium. Lee et al.75 polymerized methyl methacrylate (MMA), styrene and t-­butyl acrylate (tBA) from the GO surface using the ATRP technique. They coupled BIB on the GO surface by esterification between BIB and the hydroxyl group of GO and then polymerized the monomer from the macroinitiator on the GO surface by the ATRP technique. The modified GO exhibits a homogeneous and stable dispersion in toluene, DMF, dichloromethane and chloroform. They have tuned the chain length of grafted PS by controlling the ratio of macroinitiator and monomer, and the solution of PS-­g-­GO with the shortest polymer chains exhibits the darkest colour as it contains highest amount of GO. To make a compatible blend with poly(vinylidene fluoride) (PVDF) Layek et al.76 synthesized GO grafted PMMA by anchoring BIB on the hydroxyl groups of GO, following polymerization with methyl acrylate (MA) and subsequent reduction of graphene grafted polymer by hydrazine (Figure 2.24). The rGO-­ g-­PMMA (MG) shows a good dispersion in the PVDF matrix; this composite exhibits an increases in storage modulus, stress at break and Young's modulus of 124%, 157% and 321%, respectively, from that of pure PVDF for 5% MG. In another work, GO is linked with ethylene glycol via the esterification reaction to obtain a large number of hydroxyl groups on the surface of GO followed by attachment of BIB, and then the GO is grafted with PMMA by the ATRP technique to produce GO-­g-­PMMA, which is highly dispersible in chloroform and acts as good nanofiller for producing a high performance nanocomposite with PMMA.77 To graft poly(dimethyl aminoethyl methacrylate) (PDMA) on GO, Yang et al.78 introduced an amine group on the GO surface by amidation reaction between 1,3 diaminopropane and the carboxylic acid groups of GO. Then ATRP initiator 2-­bromo-­2-­methylpropionyl bromide is grafted from the amine and hydroxyl groups of GO and PDMA is grafted from it by in situ ATRP of the monomer. The GO-­g-­PDMA exhibits good solubility in short chain alcohol and in acidic aqueous solution. This is used to make a functional polymer–graphene composite with poly(EGDMA-­co-­ MAA) via hydrogen bonding between amine groups of PDMA and carboxylic acid group of MAA units of the copolymer. Li et al.79 used the SI-­ATRP technique to produce solution-­processable poly(tert-­butyl acrylate)-­grafted

Covalent Functionalization of Polymers

51

Figure 2.24  Synthesis  of surface functionalized graphene oxide via attachment

of ATRP initiator followed by polymerization of methyl methacrylate. (MMA = methyl methacrylate; PMDETA = N,N,N′,N″pentamethyldiethylenetriamine). Reproduced from ref. 76 with permission from Elsevier, Copyright 2010.

GO (GO-­g-­PtBA) nanosheets, which substantially enhance its dispersion in organic solvents. Thus they made a poly(3-­hexylthiophene)(P3HT)/GO-­g-­ PtBA composite thin film that exhibits nonvolatile electronic memory and bistable electrical conductivity switching behaviour with 5 wt% GO-­g-­PtBA. On hydrolysis, GO-­g-­PtBA produces water-­dispersible GO-­g-­poly(acrylic acid) (GO-­g-­PAA) allowing the decoration of Au NPs on the nanocomposite surface. For increased hydroxyl group density, Wang et al.80 produced hydroxy modified graphene by diazonium coupling between rGO and 2-­(4-­aminophenyl) ethanol. Then by esterification reaction between the hydroxyl group of hydroxy modified graphene and 2-­bromoisobutyl bromide, the ATRP initiator is covalently attached. Then using SI-­ATRP poly(2-­(ethyl phenyl amino) ethyl methacrylate) (PEMA) is grafted from the GO surface to get (G-­PEMA). Finally, they introduced the azo chromophores via azo-­coupling of G-­PEMA with the diazonium salt of 4-­aminobenzonitrile on the side-­chains of the polymer exhibiting photo-­responsive properties. Chang et al.81 synthesized a molecularly imprinted GO/poly(methacrylamide) (PMAAM) hybrid (GO–MIP)

52

Chapter 2

by transforming the carboxylic acid group of GO to acid chloride and then anchored ATRP initiator on the GO surface by reaction of acid chloride of GO and 2-­hydroxylethyl-­20-­bromoisobutyrate. Finally the PMAAM has been grafted from the GO surface using the ATRP technique. This GO–MIP is used to selectively detect 2,4-­dichlorophenol (2,4-­DCP) because the higher affinity of the GO–MIP for 2,4-­DCP over the others is for the electron transfer from the π orbitals of graphene to 2,4-­DCP causing π–π stacking. The synthesis of different industrially attractive macromolecular architectures, e.g. block copolymers, ionomers, star, dendrimers, hyperbranched or thermoplastic elastomers, can be achieved by the ATRP technique. The advent of ATRP initiators for continuous activator regeneration (ICAR),82 single-­electron-­transfer living radical polymerization (SET-­LRP),83 activator generated by electron transfer (AGET),84 activator re-­generated by electron transfer (ARGET)85 or electrochemically mediated ATRP (eATRP)86 restricts the use of the amount of copper catalyst to the few ppm level without trailing the control over polymerization making the ATRP process more smart commercially. In SET-­LRP the use of Cu(0) as ATRP catalyst has received significant attention due to its high polymerization rate at low temperature.87 It is anticipated that disproportionation of Cu(i) occurs in suitable solvents producing Cu(0) and Cu(ii). In contrast to the activation in ATRP by Cu(i), the activation takes place by reaction of alkyl halide with Cu(0), and deactivation happens by reaction with Cu(ii) as it occurs in ATRP.88 Deng et al.89 grafted GO with poly[poly(ethylene glycol) ethyl ether methacrylate] (PPEGEEMA) by grafting using SET-­LRP procedure (Figure 2.25).

Figure 2.25  In  situ growing of PDEGEEMA polymer chains via SET-­LRP from the surface of TRIS modified GO. Reproduced from ref. 89 with permission from John Wiley & Sons, Copyright 2011 Wiley Periodicals, Inc.

Covalent Functionalization of Polymers

53

To augment the number of hydroxyl groups of GO, its epoxide groups react with the amine groups of tris(hydroxymethyl) aminomethane (TRIS), then alkyl halide groups are attached by esterification reaction between hydroxyl groups and 2-­bromoisobutyryl bromide. SET-­LRP of PEGEEMA was made using CuBr/tris(2-­dimethylamino)ethyl)amine) (Me6TREN) in water/THF medium at 40 °C. Jiang et al. used a Cu(i)-­ATRP process where Cu(i) disproportionates rapidly into Cu(0) and Cu(ii) in the solvent medium before activation and polymerization can occur via Cu (0).90 Covalent linkage between PPEGEEMA and GO is confirmed by 1H NMR and FTIR data. This TRIS-­GO-­PPEGEEMA hybrid material shows reversible self-­assembly and de-­assembly in water by switching at 34 °C and it promises important potential applications in thermally responsive nanodevices and microfluidic switches. To prepare rGO sheets grafted with poly(tertbutyl methacrylate) (PtBMA), Chen et al.91 used SET-­LRP in a grafting-­from approach. In this work, the rGO is reacted with 2-­(4-­aminophenyl) ethanol and isoamyl nitrite mixture to introduce additional hydroxyl functionalities that are functionalized with bromopropionyl bromide to introduce a SET-­LRP initiation site. SET-­LRP of tBMA in DMSO in the presence of the initiator 2-­bromopropionate was conducted using Cu(0) wire and Me6TREN at 25 °C for 24 h. The 1H NMR and Raman spectra confirm the successful covalent linkage of poly(tBMA) to the GO sheets. The modified graphene nanosheets can be successfully dispersed in various organic solvents after simple sonication. The presence of the tertiary amine during SI-­ATRP offers dual functionality, e.g. reduction of GO and progress of polymer chain growth simultaneously.92 By varying the amount of ligand and GO, the reduction of GO can be controlled, hence tuning of conductivity can be made. Such a method has been utilized in the preparation of GO-­PMMA dispersion in silicone oil and their electrorheological (ER) behaviour has been studied. These suspensions are smart materials, changing their rheological properties upon an external electric field. The present GO-­PMMA system is very effective and at low particle concentrations (10 wt%) the shear stresses increases from 5 to 95 Pa, which depends on external electric field and conductivity of the particles. Singha and his coworkers93 modified GO by growing a poly(methacryloisobutyl POSS) (PMAPOSS) brush on GO surface via SI-­ATRP. GO-­PMAPOSS based semi-­IPN hydrogel was prepared by in situ copolymerization of acrylamide in the presence of soluble starch in aqueous GO-­PMAPOSS suspension. This composite hydrogel exhibits noteworthy improvement in hemocompatibility, as evident from a hemolysis test, platelet adhesion test and cytotoxicity assay. Also, this composite hydrogel is capable of controlled release of antibiotic drug Ciprofloxacin. Poly(poly (ethylene glycol) methyl ether methacrylate) (P(PEGMA)) is grafted from GO by making amino-­functionalized GO (GO-­NH2) from the reaction of surface hydroxyl groups of GO with aminopropyltriethoxysilane (APTES) followed by reacting with 2-­bromoisobutyrate to produce GO-­Br.94 The GO-­g-­P(PEGMA) is then synthesized by SI-­ATRP using CuBr/DMETA in aqueous medium at 40 °C under nitrogen atmosphere (Figure 2.26).

54

Chapter 2

Figure 2.26  The  synthesis route of GO-­g-­P(PEGMA) nanofillers via SI-­ATRP. Reproduced from ref. 94 with permission from Elsevier, Copyright 2017.

A novel nanocomposite membrane is prepared via a non-­solvent induced phase separation (NIPS)95 using polysulfone (PSF) as matrix material, NMP as solvent and GO-­g-­P(PEGMA) as filler. The new nanofillers become well dispersed in the PSF matrix, and the nanocomposite membrane shows noteworthy developments in water flux and flux recovery rate. Also, the nanocomposite membrane shows better resistance to irreversible fouling. The exceptional filtration and antifouling performance are accredited to the segregation of GO-­g-­P(PEMGA) nanofillers at the membrane surface and the nanofillers exhibit a stable retention in the nanocomposite membrane even after 30 days washing time. In a similar approach, Mahmoudian et al.96 reacted GO with 2-­bromopropinyl bromide to produce GO-­Br and then grafted MMA using CuBr/penta methyldiethylenetriamine (PMDETA) catalyst ligand system in NMP medium in the presence of a small amount of CuBr2. The GO-­PMMA is then hydrolysed in order to introduce carboxylic hydrophilic functional groups into its surface followed by impregnation into a polyethersulfone (PES) matrix to produce PES/GO-­PMMAhyd mixed matrix membranes. The incorporation of GO-­PMMAhyd into polymer matrix enhanced the hydrophilicity and water flux of the membrane. Due to the presence of acidic group of PMMAhyd in the membrane structure, pH sensitivity is noticed and a remarkable improvement in the separation of salt, heavy metal and dye rejection has been observed at alkaline pH ranges. Ohno et al.97 recently used a new ATRP initiator bearing a primary amino group, 1-­amino-­2-­((3-­((2-­bromo-­2-­methylpropanoyl)oxy)propyl)thio)-­ethane hydrochloride (BPTA), by reacting with the epoxy groups on the GO surface. They successfully grafted different molecular weight PMMA onto the GO surface using surface initiated (SI)-­ATRP and these polymer-­brush-­decorated GOs (PB-­GOs) are extremely dispersible and have never been observed for PB-­GOs made by conventional methods. These PB-­GOs form lyotropic liquid

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crystals and liquid crystal formation largely depends on the graft length of the polymer. From the selection of an appropriate graft polymer, lyotropic liquid crystals are fabricated even in a low-­polar organic solvent like hexane, and also in nonvolatile ionic liquid, and in the latter medium structurally coloured materials are observed, the colour of which is easily controlled by controlling concentration. Considering the excellent dispersibility of PB-­GO and the versatility of the methodology developed herein, this study significantly improves the applied science and technology of GO, and new applications can be developed as a consequence. To prepare polymer functionalized graphene requires harsh chemical attacks, which damage the properties of pristine graphene seriously. Huang and his coworkers98 developed a novel process to prepare diarylcarbene with ATRP initiator groups, and ATRP initiators are introduced to pristine graphene via a one-­step carbene reaction. This is used as initiator to polymerize poly(2-­hydroxyethyl methacrylate) (PHEMA) by the ATRP technique to produce PHEMA-­G. Grafting improves the solubility, dispersity, interfacial adhesion, and even reactivity of graphene to produce composites with poly(p-­phenylene benzobisoxazole) (PBO). This material could then be readily blended with PBO via in situ condensation polymerization. The hydroxyl groups of HEMA forms covalent bond between the PBO and the PHEMA-­G. The typical procedure for preparing graphene/PBO composite fibres is presented in Figure 2.27. The nanocomposite fibres at very low PHEMA-­G loading (1.0 wt%) showed a tensile strength of 3.22 GPa and Young’s modulus of 139.3 GPa, which are 51.2% and 33.7% higher than pristine PBO exhibiting an excellent reinforcing efficiency.

Figure 2.27  Nondestructive  modification of graphene via surface-­initiated ATRP and the procedure for composite fibre fabrication. Reproduced from ref. 98 with permission from American Chemical Society, Copyright 2017.

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Sha et al. used surface-­initiated activators regenerated by the electron transfer atom transfer radical polymerization SI(ARGET-­ATRP) technique to graft poly(2-­hydroxyethyl methacrylate), (HEMA) onto graphene oxide nanosheets with CuBr2/PMDETA as the catalyst/ligand system with tin(ii) 2-­ethylhexanoate as reducing agent. Uniformly distributed poly(HEMA) brushes are anchored without trace amounts of residual Cu(ii) complex on graphene substrates. The graphene brush has good potential to interfacially interact between graphene nanosheets and proteins and also exhibit good cell viability. Zhang et al.100 dispersed Fe3O4 nanoparticles on the GO surface by the solvothermal method, and a pyrene-­terminated ATRP initiator was anchored on the GO@Fe3O4 surface by π–π interaction. Finally, SI-­ATRP of 4-­vinylphenylboronic acid (VPBA) is made from the modified GO-­Br initiator to get GO@PVPBA. The GO@PVPBA nanocomposites have selective recognition capability of glycoproteins from a system consisting of standard proteins ovalbumin (OVA), transferrin (Trf), bovine serum albumin (BSA) and lysozyme (Lyz) and exhibits a high binding capacity up to 514.8 and 445.9 mg g−1 for OVA and Trf, respectively, facilitating the capture of glycoproteins from the egg white samples directly. Fan et al.101 grafted polyacrylonitrile by SI-­ATRP by functionalizing GO with diethylenetriamine (DETA) to get GO@ DETA which is treated with 2-­bromoisobutyryl bromide to get GO@BiBB followed by polymerization with acrylonitrile (AN) using a CuBr2/PMDETA catalyst/ligand system in the presence of an ascorbic acid reductant in the DMF medium. Using 2,4,6-­triphenylpyrylium tetrafluoroborate (TPPT) photocatalyst the nitrile group of GO@PAN can be transformed into pyridine and finally added into polyacrylic acid (PAA) hydrogels formed by in situ radical polymerization. These tough hydrogels have high mechanical strength, and exhibit excellent self-­healing performance. Peng et al.102 used the reactivity between the active species of atom transfer radical addition (ATRA) and the unsaturated groups of GO for sequential SI-­ATRP to graft various polymer chains and to produce various polymer architectures, such as a linear polymer, V-­shape block polymer, multibonded polymer layer and hierarchical brush-­on-­layer polymer. The organo modification of GO is made through the ATRA reactions by the ‘grafting to’ approach. GO has been modified with 1-­bromoethylbenzene (BEB), polystyrene bromide (PS-­Br), and PVDF using CuBr as a catalyst and bipyridyl (BPY) as a ligand at 80 °C for 24 h to produce GO–BEB, GO–PS and GO–PVDF. The organo modified GOs are used as initiators to perform the SI-­ATRP to incorporate another polymer chain to the GO surfaces. The monomers used in the SI-­ATRP for GO–BEB, GO–PS, and GO–PVDF are glycidyl methacrylate (GMA), sulfobetaine methacrylate (SBMA) and N-­isopropylacrylamide (NIPAAm), respectively. The monomers are polymerized from the respective macro initiators with the help of CuBr/BPY, catalyst/ligand system at 80 °C for 24 h under nitrogen atmosphere. The resulting products are GO–BEB– PGMA, GO–PS–PSBMA and GO–PVDF–PNIPAAm, and the reaction scheme is presented in Figure 2.28.

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Figure 2.28  (a)  Functionalization of GO sheets through the atom-­transfer addition

reaction (ATRA) and (b) building up various polymer architectures on GO surfaces through ATRA and sequential surface-­initiated atom transfer radical polymerization. Reproduced from ref. 102 with permission from John Wiley & Sons, Copyright 2014 Wiley Periodicals, Inc.

Thus various polymer architectures, e.g. linear, V-­shape, and hierarchical brush-­on-­layer polymers, can be built on GO surfaces. The grafts of polymer chains on the GO surfaces enhance their organo-­compatibility with a dispersion stability in different organic solvents and polymers. Maity et al.103 also grafted stimuli responsive block copolymer from GO surface via a combination of ring opening polymerization (ROP) and SI-­ATRP. Utilizing the –OH group of GO they polymerized caprolactone by ROP using Sn(oct)2 catalyst and from the pendent –OH group of the anchored caprolactone chain they anchored the 2-­bromo isobutyryl bromide. From this macroinitiator they polymerized dimethyl aminoethyl methacrylate (DMAEMA) to produce GO-­g-­ polycaprolactone-­block-­PDMAEMA. This block copolymer shows interesting morphology, and optical and electronic properties which vary with temperature and medium pH. The concept of localized p-­and n-­doping of graphene by varying pH effecting the electrical properties is first noticed in this system.

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2.3.2  R  eversible Addition Fragmentation Chain Transfer (RAFT) Polymerization Reversible addition fragmentation chain transfer polymerization (RAFT) is a controlled living radical polymerization being associated with a series of reversible addition-­fragmentation steps based on the degenerative chain transfer as a means of converting dormant chains to active propagating radicals.104,105 The RAFT mechanism is depicted in Figure 2.29. In step I, a radical initiator produces radical species; in step II, activation of monomer occurs producing the chain radicals; in step III, the chain radicals add to the RAFT agent which is a chain transfer agent (CTA) having thiocarbonylthio group and establish an equilibrium between active and dormant species (steps III and V). The chain transfer steps which make the basis of the RAFT mechanism are degenerate as they involve a reversible transfer of the functional chain end-­group (typically a thiocarbonylthio group, Z–C(=S)S–R) between the dormant chains (macro RAFT agent or macro CTA) and the propagating radicals. In an efficient process, the rate of addition/fragmentation equilibrium is much higher than that of the propagation, so there would be less than one monomer unit added per activation cycle (step V). Thus, all chains have almost same degree of polymerization (DP) at a given time. The overall process is comprised

Figure 2.29  Proposed  mechanism of reversible addition-­fragmentation chain transfer polymerization. Reproduced from ref. 105, https://pubs.acs. org/doi/abs/10.1021/acs.macromol.7b00767, with permission from American Chemical Society, Copyright 2017. Further permissions requests related to the material excerpted should be directed to the ACS.

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of the insertion of monomers between the R– and Z–C(=S)S groups of a RAFT agent, which form the α and ω end-­group of the majority of the resulting polymeric chains. In step VI termination of some propagating or initiating radical occurs; however, the majority of polymer chains remain dormant, facilitating further polymerization to produce homo or block copolymers. The RAFT polymerization depends on the high chain transfer coefficients of thiocarbonylthio compounds and trithiocarbonates, making it useful in the polymer functionalization of graphene from the grafting-­from approach.106–110 Yang et al. synthesised a RAFT reagent containing and azide group and functionalized rGO by the click reaction followed by polymerization of NIPAm by RAFT polymerization technique (Figure 2.30). For this purpose an alkyne derivative of rGO is synthesized by diazotization of aryl diazonium salts containing alkyne groups with rGO followed by the esterification reaction between 3-­azido-­1-­propanol and S-­1-­dodecyl-­S′-­(α,α′-­dimethyl-­α″-­acetic acid) trithiocarbonate producing azido-­terminated RAFT reagent. The latter is immobilized on an alkyne derivative of rGO using click chemistry and an azido-­terminated RAFT chain transfer agent (RAFT-­CTA) modified RGO is dispersed in DMF under ultrasonication. Finally NIPAm monomer and AIBN are added into the dispersion and RAFT polymerization is conducted under nitrogen atmosphere at 60 °C. This PNIPAm functionalized rGO shows a LCST transition at 33.2 °C which is lower than that of PNIPAm because of the presence of hydrophobic rGO group. The LCST measured by the heating process is almost reversible in the cooling process (LCST = 31.4 °C)

Figure 2.30  Outline  for the preparation of PNIPAM/reduced graphene oxide (RGO)

nanocomposites based on click chemistry and RAFT polymerization. Reproduced from ref. 106 with permission from John Wiley & Sons, Copyright 2011 Wiley Periodicals, Inc.

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and these modified graphene aggregates can be redispersed by cooling down below the LCST. Zhang et al.107 grafted poly(N-­vinyl carbazole) (PVK) from the GO surface to produce GO-­PVK by esterification between hydroxyl group of GO and S-­1-­dodecyl-­S′-­(α,α′-­dimethyl-­α″-­acetic acid) trithiocarbonate (DDAT) to form GO-­DDAT, which acts as a RAFT agent for the synthesis of PVK functionalized GO (GO-­PVK) (Figure 2.31). The resulting GO-­PVK hybrid exhibits a superior solubility in organic solvents and shows bistable electrical switching and nonvolatile rewritable memory effect. The RAFT polymerization technique is also used to prepare a molecularly imprinted GO/poly(methacrylamide) (PMAAm) (GO–MIP) hybrid.108 Here, the carboxylic acid group of GO is transformed into acid chloride followed by attachment with 2-­hydroxylethyl-­20-­bromoisobutyrate forming GO-­Br. The dithioester is anchored to this GO-­Br by bromide replacement with PhC(S) SMgBr. The resulting GO–MIP hybrid in aqueous solution shows a specific affinity towards 2,4-­dichlorophenol (2,4-­DCP) among the structurally related compounds due to the π–π stacking interaction. Beckert et al.109 anchored polystyrene brushes from graphene surface using sulfur-­functionalized graphene (S-­FG) as a macro chain transfer agent by means of RAFT-­mediated polymerization taking dithiourethane-­, dithioester-­ and dithiocarbonate-­ functionalized graphene. The RAFT agents are made by deprotonation of FG hydroxyl groups by a strong base and reaction with CS2 followed by alkylation. The polystyrene-­graphene brushes self-­assemble forming nanoribbons and skeleton-­like carbon superstructures during the melt processing. The negatively charged poly(acrylic acid), positively charged poly(dimethyl aminoethyl acrylate) and neutral polystyrene are successfully grafted from graphene surface by the ‘grafting from’ technique using the RAFT polymerization procedure.110 Here a pyrene functionalized graphene RAFT agent (PFRA) produced

Figure 2.31  Synthesis  of GO-­DDAT and PVK-­GO (NVC = n-­vinylcarbazole). Reproduced from ref. 107 with permission from John Wiley & Sons, Copyright 2011 Wiley Periodicals, Inc.

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by π–π stacking between a pyrene terminal RAFT agent and graphene is used for the RAFT polymerization.110 The pristine graphene exhibits ten times higher conductivity from that of different polymer grafted graphene and the different polymer grafted graphene exhibits similar conductivity to the same degree of polymerization. The longer the chain length of graphene grafted polymer the lower the conductivity. Ye et al.27 grafted graphene with PMMA, PS, PDMA, PMAA and P4VP using the RAFT technique, producing polymer brushes with multifunctional arms exhibiting solubility in water, oil, polar and apolar solvents. Graphene is functionalized with RAFT agent by introduction of alkyne groups on the surface of graphene using diazonium chemistry followed by click chemistry between an azido terminated RAFT reagent and alkyne derivative of graphene. The solubility of polymer-­functionalized graphene (FG) is strongly dependent on the nature of the grafted polymer, and it can be adjusted from water-­soluble to oil-­soluble, polar to apolar and acidic to basic by selecting a suitable polymer for grafting. Ding et al. grafted polystyrene (PS) chains on the surface of thermally reduced graphene oxide (TRGO) layers by RAFT polymerization. In this procedure, a RAFT agent, 4-­Cyano-­4-­[(dodecylsulfanylthiocarbonyl) sulfanyl] pentanoic acid, was anchored on the TRGO surface to obtain the precursor TRGO-­RAFT. Styrene polymerization was conducted using AIBN as the primary radical source and TRGO-­RAFT as the chain transfer agent. The typical procedure to prepare PS/graphene (PRG) composites is presented in Figure 2.32.

Figure 2.32  The  overall synthesis routes for the preparation of PRG composites via RAFT polymerization. Reproduced from ref. 111 with permission from Elsevier, Copyright 2014.

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The grafting density of PS/graphene (PRG) composites was measured to be 0.18 PS chains per 100 carbons of graphene. The thermal conductivity improved from 0.15 W m−1 K−1 of neat PS to 0.25 W m−1 K−1 of PRG composites with a loading of 10 wt% graphene sheets. Possibly the increase in thermal properties of PRG composites is due to the homogeneous dispersion of graphene sheets in the PS matrix forming the PRG composites. Roghani-­Mamaqani and his co-­worker112 had grafted polystyrene by a grafting from approach by using (3-­aminopropyl) triethoxysilane (APTES) as the coupling agent between hydroxyl groups of GO and the carboxylic groups of 2-­(dodecylthiocarbonothioylthio)-­2-­methylpropionic acid (RAFT agent, RA) to prepare RA-­functionalized graphene layers. Then, from the surface-­anchored RA styrene is polymerized using AIBN as primary initiator in various graft densities. Electron microscopic images show that flat layers of graphite change to a smooth wrinkled surface after oxidation and turns to opaque layers by grafting PS. Layek et al.113 grafted poly (N-­vinyl pyrrolidone) from the GO surface using RAFT polymerization where GO is first functionalized with the RAFT agent O-­ethyl xanthate via the hydroxyl groups of GO. Polymerization of N-­vinyl pyrrolidone is made using the GO-­RAFT agent in DMF at 80 °C with AIBN as primary initiator. The GO-­PVP is amphiphilic in nature and acts as a good reinforcement agent of poly(vinyl acetate) in the composite. In a similar process Kochameshki et al.114 grafted poly(diallyldimethylammonium chloride) (PDADMAC) via RAFT polymerization using xanthates as efficient chain transfer agent which are anchored on the GO sheet by 2-­bromopropionyl bromide as linking agent. The GO-­PDADMAC is used in producing nanocomposite polysulfone membranes. These modified membrane exhibits better hydrophilicity, water flux, water uptake and antifouling properties. The membrane shows the properties of heavy metal ion (Cd2+ and Cu2+) and dye methylene blue (MB) and methyl orange (MO) removal properties due to the zwitterionic nature of GO-­PDADMAC. Zhao et al.115 grafted 1,4-­divinylbenzene and 4-­vinylbenzyl chloride as monomers in water from the GO sheet by RAFT polymerization and for this GO is functionalized with 1,3-­diaminopropane in the presence N-­hydroxy-­succinimide (NHS), 1-­(3-­dimethylaminopropyl)-­3-­ ethyl­carbodiimide hydrochloride (EDC·HCl), and then by acylchloride-­2-­ (dodecylthiocarbonothioylthio)-­2-­methylpropionic acid (DDAT), which is a widely used as highly soluble chain transfer agent (CTA) for RAFT polymerization, in the presence of Et3N. Using GO-­DDAT as a 2D CTA, RAFT emulsion polymerization is carried out on graphene surfaces using 1,4-­divinylbenzene and 4-­vinylbenzyl chloride as monomers in water along with AIBN, poly(vinyl alcohol) (PVA) and NaCl. Then, the copolymer-­modified GO (GO-­PVD) produced is hyper-­cross-­linked by using FeCl3-­catalysed Friedel–Crafts reaction. Finally, graphene sandwich-­hyper-­cross-­linked microporous polymers (GHCPs) are then obtained by vacuum drying. The complete polymerization and hybrid formation process is summarized in Figure 2.33. Here, GO-­DDAT exhibits very good dispersibility in organic solvents and it acts as a 2D template for polymerization. The uniform colour of the GHCPs

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Figure 2.33  Preparation  of GHCPs: (i) N-­hydroxy-­succinimide (NHS), 1-­ethyl-­3

-­(3-­dimethylaminopropyl)carbodiimide hydrochloride (EDC·HCl), 1,3-­diaminopropane, water, 0 °C, 12 h; (ii) N2, acyl chloride-­S-­1-­ dodecyl-­S′-­(α,α′-­dimethyl-­α″-­acetic acid) trithiocarbonate (DDAT), dry dimethylformamide, Et3N, 0 °C, 24 h; (iii) N2, water, polyvinyl alcohol (PVA), NaCl, AIBN, 1,4-­divinylbenzene, 4-­vinylbenzyl chloride, 80 °C, 8 h; (iv) 1,2-­dichloroethane, FeCl3, 80 °C, 12 h. Reproduced from ref. 115 with permission from the Royal Society of Chemistry.

indicates that the copolymer becomes grafted uniformly to the GO flakes. The unadorned hyper-­cross-­linked porous polymer (HCP) was also prepared without using the graphene template for comparison. The resulting GHCPs exhibited superhydrophobic behaviour, and had a predominance of micropores with a specific surface area of up to 1224 m2 g−1. Moreover, they showed improved thermal stability in comparison with unadorned HCPs without GO. The enhanced H2 storage capacity (1.27 wt%) at 77 K and CO2 storage capacity (9.9 wt%) at 273 K at a relative pressure of 0.99 in GHPCs compared to the unadorned HCP without using the graphene template, together with high specific surface area, may be attributed to the contribution of the 2D graphene template. Han et al.116 recently synthesized, two core–shell structured reduced graphene oxide (rGO)@poly{5-­bis[(4-­trifluoro-­methoxyphenyl) oxycarbonyl]styrene} (PTFMS) with different shell thicknesses using an SI-­ RAFT polymerization technique, and are denoted as rGO@PTFMS-­1 with a thin shell and rGO@PTFMS-­2 with a thick shell where the shell thickness is tailored by the degree of polymerization of PTFMS. The effect of interfacial thickness on the dielectric behaviour of the rGO@PTFMS/poly(vinylidene fluoride–trifluoroethylene–chlorotrifluoroethylene) P(VDF–TrFE–CTFE) nanocomposites was studied, indicating that the percolation threshold of the nanocomposites increased from 0.68 to 1.69 vol% with an increase in shell thickness. Compared to the rGO@PTFMS-­1/P(VDF–TrFE–CTFE) composites,

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the rGO@PTFMS-­2/P(VDF–TrFE–CTFE) composites exhibit a higher breakdown strength and a lower dielectric constant. The results demonstrated that the grafting of rGO@PTFMS effectively enhances the dispersion and compatibility of the rGO nanofillers in the P(VDF–TrFE–CTFE) matrix. Recently Wang et al.117 grafted poly(β-­cyclodextrin)-­conjugated magnetic graphene oxide (MGO@poly(β-­CD)) for synergetic adsorption of heavy metal ions and organic pollutants. The synthesis procedure of MGO@poly(β-­CD) is illustrated in Figure 2.34. Firstly, magnetic graphene oxide (MGO) is prepared by decorating GO with magnetic Fe microspheres through a hydrothermal reaction between GO and FeCl3. Subsequently, the RAFT agent 4-­c yano-­4-­(phenylcarbonothioylthio) pentanoic acid (CPPA) is immobilized on the surface of MGO by an esterification reaction. Methyl acrylate (MA)-­β-­CD

Figure 2.34  Illustration  for the synthesis procedure of MGO@poly(β-­CD). Reproduced from ref. 117 with permission from the Royal Society of Chemistry.

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is prepared by the reaction of ethylelene diamine (EDA) modified β-­CD with glycidyl methacrylate (GMA) at 60 °C in DMF for 6 h. MGO@poly(β-­CD) is prepared by polymerizing MA-­β-­CD in MGO-­CPPA in the presence of AIBN in DMF medium under N2 atmosphere at 65 °C for 12 h and the residue was washed with DMF followed by drying. The adsorption behaviours of Cd2+ and sulfamethazine (SMT) by MGO@poly(b-­CD) in the single and binary systems indicate that the adsorption amounts of both Cd2+ and SMT are pH and ionic-­strength dependent. The sorption of Cd2+ is due to complexation and electrostatic interactions, while the removal of SMT is due to hydrogen bonding, host–guest supramolecular and π–π interactions. The synergetic adsorption of Cd2+ and SMT in the binary systems is due to interactions between loaded Cd2+ and free SMT in solutions. Namvari et al.118 used a RAFT chain transfer agent (RAFT-­CTA, S,S-­Bis(-­dimethyl-­acetic acid)-­trithiocarbonate (BDATC))-­modified rGO nanosheets (CTA-­rGONSs) by crosslinking rGONSs with a RAFT-­CTA via an esterification reaction. These nano CTA-­rGONSs are used to polymerize a hydrophobic amino acid-­based methacrylamide (N-­ acryloyl-­l-­phenylalanine methyl ester) monomer with different monomer/ initiator ratios. TGA results exhibit that the polymer–graphene composites are thermally more stable than GO itself and Mn of the polymers increases with increasing monomer/initiator ratio, whereas the polydispersity index decreases, indicating controlled polymerization. The hybrid form stable dispersion in DMF even after two months.

2.4  ‘Grafting to’ Versus ‘Grafting from’ Technique Both the ‘grafting to’ and ‘grafting from’ techniques afford control over the grafted polymer from its chain length, molecular weight, etc. In the ‘grafting-­to’ technique covalent binding of pre-­synthesized polymer chains with chain end functionalization takes place to the surface of complementary functionalized GO sheets at suitable conditions. The advantage of pre-­ synthesized polymer is that it can be synthesized with the help of the controlled living polymerization technique from outside. However, during its attachment to the functionalized GO surface there is a probability of low graft density arising, mostly from steric hindrance. So, it is very difficult to control the graft density of grafted polymer chains in the ‘grafting to’ technique. Additionally, the anchoring of a pre-­synthesized polymer chain on the graphene surface usually requires a long reaction time for the low diffusion constant of the polymer. On the other hand, in the ‘grafting from’ approach the polymerization of monomers occurs starting from the GO surface being previously modified with a suitable initiator moiety. The rate controlling step of propagation depends on the diffusion of monomers to the chain ends that are growing from the graphene surface, producing a distinct brush-­like structure with high grafting density. Thus the ‘grafting from’ approach surmounts the low grafting density and slow reactivity problem associated with the ‘grafting to’ approach. Another benefit of the ‘grafting from’ approach is the opportunity

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to produce thinner graphene sheets as the initiation begins from the initiator anchored at the graphene surface by attaching the monomer followed by propagation. During the propagation process the interlayer spacing of graphene may gradually enlarge with increasing the size of the growing chain causing generation of thin graphene sheets for the detachment from the stacked graphene sheet. In the ‘grafting to’ approach this is a difficult proposal. Hence, the ‘grafting from’ procedure is a better approach than the ‘grafting to’ technique for polymer functionalization of graphene particularly for the configuration of high graft density and formation of processable thin graphene sheets in a faster way.

2.5  Scope Covalent functionalization of graphene with polymer shows great prospects for preparation of solution processable graphene and hence graphene-­based polymer nanocomposites are suitable for different applications. Two different types of techniques, i.e. the ‘grafting to’ and ‘grafting from’ approaches, are thoroughly studied. This helps functionalization of graphene by well-­ defined polymers with precise control over molecular weight and polydispersity. The polymer functionalized graphene facilitate the introduction of different architectures onto the graphene surface, helping to produce tailor-­ made materials for fulfilling the demands in optoelectronic, energy and engineering materials. In the covalent functionalization the ‘grafting from’ technique is superior to the ‘grafting to’ technique as steric crowding can restrict an upper limit on the grafting density. Though the functionalization of graphene with polymer is plenty there is still immense scope for the development of tailor-­made properties by attaching block copolymers, dendrtic polymers, hyperbranched polymers, etc. To the graphene surface using ATRP and raft polymerization techniques. Special attention can be given to attaching graphene with electroactive polymers, which may find application to fabricate dielectric materials, supercapacitors, photovoltaic devices, etc. So, the main thrust of future research of covalently functionalized polymer graphene systems may be aimed at fabricating bio-­medical devices, biosensors, membranes, optoelectronic materials, semiconducting chips, photovoltaics, supercapacitors, solid state batteries, as well as the fabrication of high performance engineering polymer nanocomposites. In general, the diversity of tunable polymer structures on the graphene surface would be instrumental in promoting graphene from the laboratory to many technological and biotechnological appliances.

References 1. W. S. Hummer and R. Offeman, J. Am. Chem. Soc., 1958, 80, 1339. 2. P. Ren, D. Yan, X. Ji, T. Chen and Z. Li, Nanotechnology, 2011, 22, 055705. 3. S. Stankovich, D. A. Dikin, R. D. Piner, K. A. Kohlhaas, A. Kleinhammes, Y. Jia, Y. Wu, S. T. Nguyen and R. S. Ruoff, Carbon, 2007, 45, 1558.

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Chapter 3

Noncovalent Polymer Functionalization of Graphene 3.1  Introduction In spite of a large number of methods for covalent functionalization of polymers on graphene surface, the noncovalent functionalization of small molecules and polymers on graphene surface has also been used significantly.1–9 In noncovalent interactions, H-­bonding and π–π stacking on the graphene surface play a major role. Noncovalent functionalization possesses noteworthy advantages over covalent functionalization because it augments the solubility without varying the extended π conjugation of the graphene sheet, whereas in the covalent functionalization new sp3 defects are produced on the graphene surface decreasing the π conjugation. We shall discuss the involvement of different interactions involved in noncovalent functionalization in the following section, although sometimes it involves one or more noncovalent interactions in the same system.

3.2  π-­stacking Interaction Xu et al. produced stable water dispersions of graphene sheets using an aqueous solution of pyrene derivative, 1-­pyrene butyrate (PB), which acts as a stabilizer because the pyrene moiety has strong affinity with the basal plane of graphite via π-­stacking.1 This dispersion yields a flexible graphene film showing a conductivity of 2 × 102 S m−1. Self-­assembled monolayers on epitaxial graphene were produced by Wang et al.2 with the molecular semiconductor perylene-­3,4,9,10-­tetracarboxylic dianhydride (PTCDA) grown on a SiC (0001) surface. The hybrid possessed long-­range order as evident   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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from ultra-­high vacuum scanning tunnelling microscopy. Mullen and his coworkers used a noncovalent functionalization method to make a conjugate of graphene with an electronic donor, sodium salt of pyrene-­1-­sulfonic acid (PyS), and an electronic acceptor, disodium salt of 3,4,9,10-­perylene tetracarboxylic diimide bisbenzene sulfonic acid (PDI) via π–π interactions without disrupting the electronic conjugation of graphene.3 The aqueous dispersion of the hybrid graphene sheets are highly stable because negative charges present in both the dispersant molecules act as stabilizing agents for the presence of strong static repulsive force between the negatively charged reduced graphene sheets. They used the film of graphene conjugate as electrodes in the bulk heterojunction solar cells, which showed greatly improved power conversion efficiency. Lomeda et al. used sodium dodecylbenzenesulfonate (SDBS) surfactant to wrap rGO sheets followed by treating with aryl diazonium salts4 forming the conjugate by π-­stacking. Such functionalized graphene nanosheets disperse readily in polar aprotic solvents. Liu et al.5 noncovalently functionalized PNIPAAM to make thermosensitive graphene–polymer composites via π–π interaction between graphene and pyrene terminated PNIPAAm (Figure 3.1). To achieve this, they synthesized a pyrene-­terminated thiocarbonylthio RAFT agent which was used to polymerize NIPAAm to obtain pyrene terminated PNIPAAm to bind GO by π–π interaction. This noncovalently functionalized hybrid material gives stable homogeneous dispersion in aqueous solution and is thermosensitive showing a lower critical solution temperature (LCST) at 24 °C from that of the pristine polymer (33 °C). Thus the

Figure 3.1  A  scheme depicting the synthesis of pyrene-­terminated PNIPAAm using

a pyrene-­functional RAFT agent and the subsequent attachment of the polymer to graphene. Reproduced from ref. 5 with permission from John Wiley & Sons, Copyright 2009 Wiley Periodicals, Inc.

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property modification of the noncovalent functionalization of graphene with pyrene terminated PNIPAAm is very similar to that of covalent binding of PNIPAAm to GO. Similarly pyrene terminated poly(dimethylamino ethyl methacrylate) (PDMA) and poly(acrylic acid) (PAA) have been synthesized using RAFT polymerization and the graphene is functionalized with these polymers by π-­stacking interaction. The formation of self-­assembly of the two oppositely charged graphene–polymer composites through noncovalent interactions form layer-­by-­layer structures.6 At different pH values these hybrid materials exhibit phase transfer behaviour between organic and aqueous media. The multipyrene terminated hyperbranched polyglycidol (mPHP) has been used by Li et al.7 to noncovalently functionalized pristine graphene sheets (GSs) through π–π stacking. This modified GS can generate and stabilize a diversity of metal nanoparticles (Ag, Au and Pt) to produce different GS/mPHP/ metal nanohybrids. The GS/mPHP/Au hybrid exhibits good SERS activity facilitating its use as an efficient heterogeneous catalyst for the reduction of 4-­nitro phenol. Xu et al.8 functionalized rGO with perylene bisimide-­ containing poly(glyceryl acrylate) (PBIPGA) via π–π stacking interaction where bifunctional N,N′-­bis{2-­[2-­[(2-­bromo-­2-­methylpropanoyl) oxy]ethoxy] ethyl}perylene-­3,4,9,10-­tetracarboxylic acid bisimide (PBI–Br) is used as ATRP initiator to synthesize perylene bisimide containing poly(solketal acrylate) which on hydrolysis yields PBIPGA. The rGO-­PBIPGA conjugate exhibits good dispersity in water and has very low cytotoxicity toward 3T3 fibroblasts after 6 and 24 h of incubation. In the presence of a conducting polymer poly(3,4-­ethylene dioxythiophene) and poly(styrene sulfonate) (PEDOT:PSS) reduction of GO occurs accompanied by noncovalent functionalization of the conducting polymer with rGO.9 A strong π–π interaction between the two-­dimensional rGO and rigid backbone of PEDOT and the intermolecular electrostatic repulsions between negatively charged PSS causes the noncovalent functionalization. The resultant rGO/PEDOT:PSS hybrid exhibits high conductivity with a tuneable transmittance and its dispersion exhibits a good colloidal stability in aqueous medium. The network structure of conducting polymers provides good flexibility and mechanical stability towards rGO nanosheets facilitating its application as highly flexible and transparent electrodes. Bai et al.10 have noncovalently functionalized graphene with sulphonated polyaniline (SPANI) formed during reduction of GO with hydrazine hydrate in the presence of SPANI. The resulting SPANI/rGO hybrid forms a good dispersion in water and the composite films show enhanced electrocatalytic activity and improved electrochemical stability. An amphiphilic coil-­rod-­coil conjugated triblock copolymer (PEG-­OPE-­ PEG) is noncovalently functionalized with rGO by π–π interaction when GO is reduced with hydrazine in the presence of PEG-­OPE-­PEG (Figure 3.2).11 The PEG-­OPE-­PEG possesses one lipophilic π-­conjugated oligomer and two hydrophilic PEG coils. Hence, the hybrid PEG-­OPE-­PEG-­rGO exhibits good dispersion in both aqueous and organic solvents. Peponi et al.12 used the noncovalent π-­stacking interaction to investigate the self-­organization of

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Figure 3.2  (a)  Chemical structure of PEG-­OPE-­PEG. (b) The synthesis of PEG-­OPE-­ PEG-­rGO in H2O. Step 1: oxidation of graphite yields single-­layer GO sheets. Step 2: chemical reduction of GO with hydrazine in the presence of PEG-­OPE-­PEG produces a stable aqueous suspension of PEG-­OPE-­PEG-­ rGO. (c) Photograph of (A) GO and (B) rGO in water, (C) PEG-­OPE-­PEG-­rGO and (D) PEG-­OPE-­PEG in methanol. Reproduced from ref. 11 with permission from John Wiley & Sons, Copyright 2009 Wiley Periodicals, Inc.

the poly(styrene-­b-­isoprene-­b-­styrene) (SIS) block copolymer matrix with graphene nanosheets for fabricating transparent thin films for optoelectronic applications. The noncovalently functionalized rGO with poly (sodium 4-­styrenesulfonate) are prepared by coating rGO nanoplatelets with an amphiphilic poly(sodium 4-­styrenesulfonate) by π-­stacking interaction producing stable aqueous dispersions of graphitic nanoplatelets.13 Ji et al.14 successfully synthesized pyrene-­terminated liquid crystalline polymers (poly[8-­(4-­c yano-­40-­biphenyl)-­1-­octanoylacrylate] and poly[6-­ (4-­cyano-­40-­biphenyl)-­1-­hexanoylacrylate]) and functionalized them with graphene sheets successfully via π–π interactions (Figure 3.3). The resulting complexes exhibit excellent dispersion stability in DMF; however, they show poor dispersibility in aqueous solutions. FT-­IR, TGA and AFM results indicate that the liquid crystalline polymers (LCPs) are attached on the surface of graphene sheets through noncovalent bonds and the fluorescence quenching of pyrene anchoring unit of LCP indicates that pyrene moiety strongly interacted with the surface of graphene sheets via π–π interactions. XRD results suggest that intercalated LCP chains exhibit closer packing compared to LCP chains in the bulk indicating two-­dimensional (2D) confined environment of LCP due to the π–π interactions of graphene sheets. Chen et al.15 demonstrated the utilization of poly(2-­butylaniline) (P2BA) as noncovalent dispersant for preparation of graphene/polymer coatings (Figure 3.4). P2BA is prepared by polymerizing 2-­butylaniline in HCl solution

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Figure 3.3  A  schematic illustration of functionalization of graphene by pyrene-­ terminated LCP via noncovalent interaction. Reproduced from ref. 14 with permission from Elsevier, Copyright 2015.

Figure 3.4  Dispersion  of graphene in organic solvent using poly 2-­butyl aniline (P2BA) as dispersing agent. Reproduced from ref. 15 with permission from Elsevier, Copyright 2016.

at 0 °C with ammonium persulfate and then the product is dedoped with 0.1 M NH4OH to get P2BA in emeraldine base form. The presence of aromatic heterocyclic structures in P2BA facilitates the dispersion of graphene (G) in THF through π–π interaction. This P2BA-­G is used to make composite with epoxy resin and the composite enhances the corrosion protection performance of steel. Lian et al.16 produced graphene nanoribbons (GNRs)

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by oxidative unzipping of MWCNTs with acidic KMnO4 and have modified them using Kevlar through noncovalent π–π stacking interaction between the aromatic part of Kevlar and the graphitic surface of GNRs. The modified KGNRs become well dispersed in PVC and PMMA matrices and exhibit strong interfacial interaction between them. This improves the mechanical properties of PVC and PMMA composite films significantly, making the KGNRs ideal fillers for fabrication of high performance polymer composites. Maity et al.17 made noncovalent functionalization of rGO with the conducting graft copolymer (polythiophene-­g raft-­polymethylmethacrylate, PT-­g-­PMMA) produced by combination of oxidative polymerization of a thiophene unit followed by ATRP polymerization of methyl methacrylate. The strategy has three major steps: (i) chemical oxidation of graphite to yield GO by the Hummer technique, (ii) reduction of GO to RGO and (iii) attachment of the graft copolymer PT-­g-­PMMA to the RGO by noncovalent interactions (Figure 3.5). The dispersion ability of the polymer functionalized graphene (PFG) has largely improved as demonstrated from the digital images of PT-­g-­PMMA exhibiting a transparent yellow solution and the PFG exhibits a stable black dispersion in DMF. The noncovalent functionalization is also evident from the UV–vis spectra where the absorption band at 430 nm for π–π* transition of the conjugated thiophene chain decreases significantly on addition of rGO and the intensity of the fluorescence spectra of PT-­g-­PMMA shows significant quenching due to the efficient charge transfer through π-­stacking. PFG is used to make composites with poly(vinylidene fluoride) (PVDF), where rGOs are distributed homogenously in a PVDF matrix

Figure 3.5  Schematic  presentation of the functionalization of RGO by PT-­g-­PMMA by noncovalent interaction. Reproduced from ref. 17 with permission from Elsevier, Copyright 2016.

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and the composite with 5% PFG shows 90% piezoelectric β-­polymorph formation. The composite exhibits a maximum (60%) increase in storage modulus with the highest increment of tensile strength and strain of 317% and 302% over PVDF, respectively. Kim et al.18 recently produced dispersible graphene nanoplatelets (GNPs) using a noncovalent approach via π–π interactions with poly(4-­aminostyrene) (PAS) and the hybrid exhibits improved degree of dispersion in an epoxy matrix, likely because of a steric hindrance effect. Also, PAS acts as a molecule for maintaining the intrinsic properties of GNPs and enhancing the interfacial strength between GNPs. The PAS-­GNP/ epoxy nanocomposites had significantly higher fracture toughness (2.526 MPa m1/2 at 4 wt%) compared to that of the pure epoxy (1.01 MPa m1/2), which corresponds to an improvement of about 154%. Also, the PAS-­GNP/epoxy nanocomposites are applied to carbon fibre fabric (CF) and the interlaminar shear strength (ILSS) and fracture toughness show greatest improvements in ILSS (252%) and fracture toughness (142%) with 4 wt% PAS-­GNPs. The superior mechanical properties of the CF/PAS-­GNP/epoxy nanocomposites are attributed to the better filler dispersion and crack bridging properties. Polyketone (PK) has received significant attention as a polymer barrier material because of its outstanding performance for blocking of water vapour molecules. Nevertheless, for high-­criteria requirements in packaging applications it is necessary to improve the barrier properties of PK. Graphene nanoplatelets (GNPs) are also suitable for gas barrier materials because of their atomically two-­dimensional sheets and high aspect ratio. So, nanocomposites of PK and GNP can play an important role by blocking gas vapour diffusion with lengthening pathways for diffusing gas molecules. In order to accomplish a longer and more tortuous path for the penetrant, a good dispersion of GNP in the PK matrix is required. Cho et al.19 made PK composites with 1-­aminopyrene (APy) functionalized GNP via a noncovalent attachment using π-­stacking interaction (Figure 3.6).

Figure 3.6  Schematic  diagram for noncovalently functionalized GNP/APy and its interaction mechanism with the PK matrix. Reproduced from ref. 19 with permission from Elsevier, Copyright 2017.

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GNP/APy can effectively induce specific interaction such as hydrogen bonding between the amine functional group of GNP/APy and the carbonyl functional group of the PK matrix facilitating good dispersion of GNPs with compact interface. The homogeneous dispersion state of GNPs results in an outstanding improvement in the barrier of water vapour transmission rate. The nanocomposite film at 1 wt% loading of filler increases barrier performance by two times compared to that of PK/pristine GNP nanocomposite film, showing 92% reduction in water vapour transmission rate compared to the case of pristine PK film. Thus this facile method of graphene functionalization enhances the graphene dispersibility and interfacial interaction with the matrix polymer would be fruitful for potential applications as gas barrier materials. Song et al.20 used pyrene-­functionalized poly(methyl methacrylate)-­block-­polydimethylsiloxane (Py-­PMMA-­b-­PDMS) copolymers, synthesized via the ARGET ATRP method to functionalize GO through π–π interaction between pyrene moiety of block copolymer and GO sheets. The noncovalently functionalized GO (GO@Py-­PMMA-­b-­PDMS) is used to make composite with PMMA and, on mixing with 0.05 wt% GO@Py-­PMMA-­b-­ PDMS, the tensile strength, Young's modulus, elongation at break and toughness of PMMA increased by 23%, 54%, 117% and 218%, respectively, showing both the reinforcing and toughening effects of the functionalized GO. Chen et al.21 functionalized GO by dropping an alcoholic solution of phenoxycycloposphazene (HPTCP) into an aqueous suspension of GO when the less-­soluble HPTCP was attached to GO via strong π–π interactions, making a good dispersion of functionalized GO. Then poly(vinyl alcohol) (PVA) films were immersed in aqueous solution of polyethyleneimine (PEI) followed by rinsing and drying. The PEI-­modified PVA membranes were then immersed into the aqueous solution of functionalized GO followed by rinsing and drying. These steps were repeated until the desired number of layer-­by-­layer (LBL) deposition occurred and PVA surfaces acted as a protective shield to effectively block heat and mass transfer, thus the GO-­based LBL coating has great potential as a flame-­retardant material. Hu et al.22 functionalized rGO with bisphenol A (BPA) by refluxing a mixture of rGO dispersion and BPA solution prepared at basic pH. The BPA was adsorbed on the basal plane of rGO by π–π stacking interaction. The BET surface area of the nanocomposite was 130 m2 g−1 and that of rGO was 70 m2 g−1, indicating BPA can efficiently prevent restacking of rGO. They studied electrochemical behaviours and the results showed that the nanocomposites have specific capacitance of 466 F g−1 at a current density of 1 A g−1, having 81% retention at 10 A g−1 compared to 1 A g−1 and possess superior cycling stability. Thus, the rGO/BPA nanocomposites act as a promising electrode materials for supercapacitor applications. Hou et al.23 dispersed graphene nanosheets (GNs) with 3,4,9,10-­perylenetetracarboxylic acid anhydride (PTCDA) by π–π stacking interactions. Then in situ polymerization of ε-­caprolactone was made, producing hyperbranched polycaprolactone modified GNs (PGNs). To improve the mechanical and thermal properties of the epoxy PGNs are added into the epoxy matrix at different contents and, at 1.0 wt% PGNs content in the composite, the impact strength and tensile

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Figure 3.7  Schematic  representation of the interaction between AAPN derived

from adenine and graphene flakes. Reproduced from ref. 26 with permission from Springer Nature, Copyright 2019.

strength increased by 148% and 87%, respectively compared with those of pure epoxy, indicating the excellent reinforcing properties of PGNs. Graphene is also noncovalently functionalized with the nucleic acids DNA/ RNA for applications in biomedical fields and the π-­stacking interaction between nucleobases of nucleic acid is responsible for exfoliation and dispersion.24,25 However, Liu et al.26 recently used small biomolecules such as adenine, modified with aromatic phthalonitrile (AAPN), to functionalize graphene via π-­stacking interaction, making a stable dispersion of graphene (Figure 3.7). The interaction is judged from the UV–vis absorption spectra, where AAPN absorbed on the graphene surface is significantly red shifted (42 nm) compared to that of the pure AAPN spectrum for the absorption maxima at 307 nm. This AAPN-­G forms good and stable dispersion and has the potential for multifarious applications such as absorption for organic dyes, etc.

3.3  H-­bonding Interactions The graphene oxide has pendant –OH and –COOH groups that are capable of H-­bonding with other –OH, –COOH and –NH2 groups of polymeric chains producing noncovalently polymer functionalized graphene. Liang et al.27 used a simple water solution processing method for the preparation of poly(vinyl alcohol) (PVA) nanocomposites with GO via H-­bonding and exhibits 62% improvement in Young's modulus and 76% increase in tensile strength

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28

for the addition of 0.7 wt% GO. Layek et al. used the same water solution processing method to produce composite of PVA with sulphonated graphene (SG) instead of GO by H-­bonding interaction (Figure 3.8). The composite exhibits a change in morphology from fibrillar, dendritic to rod-­like morphology for the addition of 1, 3 and 5% (w/w) SG, respectively. The composite having dendritic morphology exhibits the highest increase of stress (177%), strain (45%) at break, toughness (657%) and storage modulus (1005%) from that of pure PVA. The same group produced nanocomposites of chitosan with SG and a highest increase of both tensile strength and Young’s modulus of 290 ± 7% and 200 ± 7%, respectively from those of chitosan for the addition of 5% SG (w/w).29 The strong noncovalent interaction involving H-­bonding has been attributed to a dramatic increase in mechanical properties in both cases. Layek et al.30 also produced high performance nanocomposites of a biodegradable and biocompatible polymer, sodium carboxymethyl cellulose (NaCMC), with graphene oxide (GO) by the solution casting method from water medium. The spectroscopic data clearly indicate the H-­bonding interaction between –COOH, –OH and the epoxy group of GO with the carboxylate ion of NaCMC is the reason for composite formation. Wang et al.31 noncovalently functionalized rGO with a triblock copolymer of poly(ethylene oxide) (PEO) and poly(propylene oxide)(PPO) (PEO-­b-­PPO-­b-­ PEO) via H-­bonding interaction to produce a mechanically strong, conducting graphene/chitosan nanocomposite. This modified graphene composite film exhibits electrical conductivity of 1.2 S m−1, Young's modulus 6.3 GPa, tensile strength 206 MPa and elongation at break 6.5% for 6% rGO content (w/w). The uniform dispersion of noncovalently functionalized rGO nanofillers in the

Figure 3.8  H-­  bonding interaction between PVA and SG. Reproduced from ref. 28 with permission from Elsevier, Copyright 2011.

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polymer matrices and formation of compact layered structure of the composite is due to the H-­bonding interaction attributing to forming the mechanically strong and electrically conducting composite films. Also the H-­bonding technique was used by Zu et al.32 to produce noncovalently functionalized rGO with PEO-­b-­PPO-­b-­PEO by in situ reduction of GO in the PEO-­b-­PPO-­b-­PEO matrix by hydrazine. This modified rGO shows stable dispersion in aqueous medium due to the hydrophilic PEO chains extending into water. The hybrid produces supramolecular hydrogel with cyclodextrin as penetration of PEO chains occurs into the cyclodextrin cavities producing well-­dispersed graphene facilitating supramolecular hydrogel formation. Choi et al.33 noncovalently functionalized rGO with amine-­terminated polystyrene due to the presence of residual carboxylic acid groups providing H-­bonding sites to the amine terminated polystyrene. On sonication this noncovalently functionalized rGO shows a stable dispersion in various organic solvents and the hybrid exhibits phase transfer of graphene sheets from the aqueous phase to the organic phase. The same technique of reduction of GO by hydrazine in presence of cationic polyelectrolyte, poly[(2-­ethyldimethylammonioethyl methacrylate ethyl sulfate)-­co-­ (1-­vinylpyrrolidone)] (PQ11), was used by Liu et al.34 to form a noncovalently functionalized rGO/PQ11 hybrid that shows a stable aqueous dispersion of graphene nanosheets. The multipoint attachment of the PQ11 chain with rGO surface by H-­bonding, and due to the electrosteric stabilization, the hybrid shows remarkable colloidal stability and the hybrid can decorate AgNPs which are able to detect H2O2 without using an enzyme. In GO the localized molecular sp2 cluster within the sp3 matrix confines the π-­electrons and the size of sp2 cluster dictates the local energy gap. So, GO on noncovalent functionalization with polymers by H-­bonds with GO can tune the cluster size, thus Kundu et al. observed an increase in the fluoresence intensity.35,36 The noncovalent attachment of GO on methyl cellulose (MC)35 (Figure 3.9) shows interesting fluorescent property of GO–MC hybrids, particularly at pH 4.

Figure 3.9  Supramolecular  interaction of the GO–MC hybrid. Reproduced from ref. 35 with permission from the Royal Society of Chemistry.

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The morphology of the GO–MC hybrids at pH 4 is ribbon type due to self-­ organization, but no characteristic morphology is observed at pH 7 and 9.2. A drastic decrease in PL intensity (91%) on addition of picric acid to the GO–MC system is noticed. Thus the hybrid system acts as a good sensor for the detection of picric acid by instantaneous photoluminescence quenching with a detectable limit of 2 ppm.35 In an acidic medium (pH 4) the GO and poly(vinyl alcohol) (PVA) hybrid produced by H-­bonding interaction exhibit good fluorescent properties.36 FTIR spectra indicate the hydrogen bond formation between hydroxy groups of GO and the hydroxy group of PVA, and Raman spectra indicate that due to passivation by H-­bonding the hybrid is highly fluorescent. The fluorescent microscopic images of the hybrids show fibrills that emit highly intensive green light. The fibrillar morphology is also noticed from FESEM micrographs and they are produced due to the supramolecular organization of GO–PVA complex. The highly fluorescent GO–PVA hybrid acts as a fascinating tool for the selective detection of Au(iii) ions in aqueous media with a detectable limit of ∼275 ppb. Pandele et al.37 functionalized GO noncovalently with a nonionic surfactant (Tergitol NP 9, f) via H-­bonding as evident from the shift of the >C=O vibration peak in the FTIR spectra due to H-­bonding interaction with the hydroxyl group of the surfactant Tergitol NP. This causes good dispersion in water and facilitates fabrication of pectin/GOf composite films via mixing in aqueous medium followed by sonication and drying. The formation of the composite is illustrated in Figure 3.10. An XRD study has indicated the intercalation of surfactant molecules between GO sheets due to hydrogen bonding and Raman spectroscopy shows an increase in defects of GO after the noncovalent modification. To make epoxy composite the PSS-­G is dispersed in alcohol and epoxy resin is added, followed by ultrasonication and drying, leading to the formation of novel hybrid materials with

Figure 3.10  Chemical  structure of pectin, Tergitol NP, schematic representation of GO and possible interactions present in pectin/GO and pectin/GOf composites. Reproduced from ref. 37 with permission from Elsevier, Copyright 2017.

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improved mechanical performance. For example, by adding 1 wt% of GOf within the polymer a significant increase in the Young’s modulus from 2.8 to 7 MPa is noticed. However, the composite shows a reduction in tensile strain from that of pristine pectin indicating an increase in the rigidity of the composite materials. Gupta et al.38 functionalized rGO nanosheets (NSs) with poly(ethylene glycol) 200 (PEG200) by mixing the component solution followed by γ-­irradiation. This method of functionalization causes intercalation of PEG200 chains in the rGO gallery through H-­bonding between the oxygen atoms of PEG200 and the hydroxyl groups of rGO molecules resulting an increase in d-­spacing of the graphene sheets. In the composite the frictional coefficient and wear resistance of sliding steel surfaces gets reduced by 38% and 55%, respectively, when 0.03 mg mL−1 PEG200-­functionalized rGO is dispersed in PEG200. The effective dispersion of PEG200 and rGO occurs arising from bipolar interactions between the components increasing the lubrication properties, and chemical analysis of wear particles exhibit decomposition of rGO into nanosized graphite domains forming effective and stable tribo films on the steel wear tracks that easily shear under contact stress. This causes significant enhancement of antifriction and antiwear properties. At high rGO concentrations, the lubrication efficiency diminishes due to graphene–graphene intersheet collisions, creating mechanical energy and chemical defects at contact boundaries.

3.4  Surfactant Induced Functionalization The noncovalent functionalization of GO with surfactant is more convenient and simpler in comparison to the other strategies, and thus it is a feasible and economic method for industrial applications of graphene. Both neutral and ionic surfactants can be used for functionalization of GO due to the presence of hydrophobic and hydrophilic groups interacting with GO by noncovalent interactions. Layek et al.39 used surfactant (pluronic F127) induced functionalization of rGO to make nanocomposites of PF127-­rGO with polyethylene glycol plasticized gum arabic (PGA). The PF127-­rGO is synthesized by refluxing aqueous dispersion of GO with hydrazine40 in the presence of PF127 when GO is reduced to rGO, making a noncovalent bond with rGO (Figure 3.11). It shows good dispersion in an aqueous medium due to the presence of the hydrophilic polyethylene oxide segment in pluronic F127. FTIR, UV–vis and Raman spectra indicate that pluronic F127 functionalization does not alter the structure of rGO, and the rGO sheets remain exfoliated in a dilute solution of PF127-­rGO, as evident from TEM images. The PF127-­rGO and PGA composite is prepared by evaporating an aqueous solution mixture of components at different concentration of the filler. The composite films characterized by FESEM and WAXS studies exhibit good dispersion of PF127-­rGO sheets in the PGA matrix. The FTIR results indicate a significant interaction

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Figure 3.11  Schematic  representation of the preparation of PF127-­rGO and PF127-­ rGO/PGA composites. Reproduced from ref. 39 with permission from Elsevier, Copyright 2017.

between functional groups of PF127-­rGO and those of PGA and PF127-­rGO 7.5 exhibits a 124% increase in stress at break and 185% increase in Young's modulus from those of pure PGA. To prepare poly(lactic acid) (PLA)/graphene oxide nanocomposites Zhang et al.41 noncovalently functionalized GO with the surfactant didodecyl dimethyl ammonium bromide (DDAB) to produce f-­GO. In this noncovalent functionalization, the hydrophilic quaternary ammonium (NR4+) heads of amphiphilic DDAB molecules are attached onto the surfaces of GO through ionic interaction with negatively charged carboxyl groups (COO–) of GO sheets, whereas the long hydrophobic C12 alkyl tails stretch toward the organic phase (Figure 3.12). So, the resulting f-­GO is easily transferred from the water phase to the chloroform phase by simple shaking. The chloroform dispersion of f-­GO is used to make a composite with PLA and the incorporation of only 0.2 wt% f-­GO into PLA results an extraordinarily (26.6 times) increase in elongation-­ at-­break compared to PLA without sacrificing tensile strengths. Such an outstanding improvement in toughness is attributed to the excellent exfoliation of GO sheets and good interfacial adhesion between PLA and modified GO nanosheets.

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Figure 3.12  Schematic  f-­GO nanocomposites together with chemical structure of DDAB molecules. Reproduced from ref. 41 with permission from Elsevier, Copyright 2017.

In a recent work, Oh et al.42 compared the effect of cationic and anionic surfactant on the formation of graphene-­coated PMMA to form composites with epoxy resin. In this process, sodium dodecyl sulfate (SDS) and cetyl triammonium bromide (CTAB) surfactants with opposite charges are used to coat the graphene surface independently with the same monomer, MMA, followed by admicellar polymerization. The resultant PMMA-­functionalized graphene, using SDS (S-­PfRG) and CTAB (C-­PfRG), showed different morphologies. S-­PfRG does not form well-­exfoliated morphology because of the anionic nature of the surfactant, which cannot anchor from the GO surface which is negatively charged and exists as an agglomerate with the PMMA chains. However, C-­PfRG exhibits thin layer morphology as the PMMA coating of GO is homogeneous due to the interaction of negatively charged GO with the cationic CTAB surfactant, making good exfoliation of GO, on which the PMMA is grown forming microspheres, thus preventing the restacking of the graphene sheets. The epoxy composites are prepared by the solution blending and casting method by dispersing the filler and epoxy-­terminated dimethylsiloxane (ETDS) homogeneously in ethanol and the composites thus produced show different mechanical and thermal properties. The C-­PfRG/epoxy composites exhibit not only well-­dispersed exfoliation but also an increase in interfacial adhesion compared to S-­PfRG/epoxy, exhibiting much

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better thermal conductivity from that of the S-­PfRG/epoxy composites. The storage modulus of the graphene/epoxy composites also improved and the C-­PfRG/epoxy composite exhibits higher modulus than S-­PfRG/epoxy composite for 1 wt% filler content.

3.5  Miscellaneous Nonbonding Interactions Very recently, much noncovalently functionalized graphene with polymer via multiple hydrophobic, π–π stacking, H-­bonding interactions have been reported. Liu et al.43 used an in situ reduction technique to produce Nafion-­rGO transparent conducting films by reducing GO/Nafion dispersion using hydrazine. In fact, the Nafion functionalized graphene is formed by hydrophobic interaction of Nafion with the graphene surface and exfoliates the graphene by an electrosteric mechanism. To obtain highly conductive graphene-­based materials noncovalent functionalization of graphene sheets is carried out with conjugated poly(2,5-­bis(3-­sulfonatopropoxy)-­1,4-­ethynylphenylene-­ alt-­1,4-­ethynylphenylene) polyelectrolyte(PPE-­SO3–).44 This resulting hybrid exhibits a high quality graphene dispersion in water due to good solubility of PPE-­SO3– in aqueous medium. Due to the attractive optoelectronic properties of PPE-­SO3– the graphene/PPE-­SO3– hybrid could realise a variety of optoelectronic applications. Kim et al.45 transformed hydrophilic GO to a hydrophobic phase transferable graphene using polymeric ionic liquid (PIL), poly(1-­vinyl-­3-­ethylimidazolium). These functionalized graphene sheets with PIL exhibit stable dispersion in the water phase which is readily transferred into the organic phase by modulating their properties from hydrophilic to hydrophobic using the anion exchange method. Layek et al.46 produced sulfonated poly(ether–ether–ketone)(SPEEK) functionalized rGO(SPG)/poly(vinylidenefluoride) (PVDF) composites by exploiting π–π and H-­bonding interactions. The benzene ring of SPEEK can interact with rGO by π–π interactions and the –SO3H group can assist in dispersing in polar solvents. The TEM study of SPG exhibits some wrinkles on the fully exfoliated graphene sheets thus helping to make composites with PVDF by dispersing in DMF, which is a common solvent of SPG and PVDF. In the composite the –SO3H group of SPG interacts with the >CF2 dipole of PVDF to disperse SPG uniformly into the PVDF matrix. A cross-­sectional FESEM micrograph exhibits folded-­chain lamella of PVDF crystallites spreading outwards from the growth centre of the spherulite. The spherulitic size of PVDF decreases with the increasing SPG content and the composite having 3 wt% SPG (SPG3) does not exhibit spherulite, but shows fibrillar crystals. FTIR and XRD results suggest that piezoelectric β-­polymorph PVDF formation starts with incorporation of SPG, and SPG3 shows the fully β-­polymorph formation. Study of the mechanical properties indicate that there is a simultaneous enhancement of strain and stress at break, indicating enhancement of toughness of the PVDF composites compared to pure PVDF. The Young's

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modulus exhibits 160% increase and oxygen permeability coefficient reduces to 91% in the SPG3 composite from the pure PVDF. Qui et al.47 have used a graphene–polydopamine system to make composite with flexible polyurethane foams (FPUFs). The neurotransmitter dopamine is a sole molecule to mimic adhesion of proteins and it undergoes self-­polymerization at dilute alkaline condition to form a highly adherent polydopamine (PDA) film on the surface of various substrates.48 The self-­polymerization of dopamine is an oxidation process in which the catechol moiety is oxidized to a quinone form and the dopamine is a green reducing agent that can not only reduce the GO, but also can closely adhere to the rGO surfaces in the form of PDA improving the dispersion and stability of rGO in water.49 PDA-­rGO is made layer-­by-­layer (LBL) assembly and, finally, it is used to construct a dense and robust flame retardant nanocoating on the surfaces of FPUFs along with two polyelectrolytes, polyacrylic acid and branched polyethyleneimine BPEI via a LbL deposition process tailored by changing pH and PDA-­rGO concentration. Cone calorimeter data indicate that a PDA-­rGO-­based coating reduces the peak heat release rate by 49.3%, and smoke release by 33.1%, compared to the neat FPUF. This improvement in flame retardant properties is attributed to the free radical scavenging property of polydopamine and the physical barrier effect of rGO sheets. Li et al.50 recently fabricated PVDF composite with exfoliated graphene sheet (EG) using π–π and H-­bonding interactions together to achieve a high dielectric constant of PVDF composites. For this purpose they modified EG with carboxyl-­enriched perylene derivative (Py) to obtain PyEG nanosheets that act as novel nanofillers. This noncovalent functionalization by π–π interaction has not only kept high conjugation of EG, but also bestowed EG with more carboxyl groups at its surface. This results in a large number of H-­bonds between F atoms of PVDF and hydroxyl groups (–OH) groups of PyEG and strong dipolar interactions between CF2 group of PVDF and >C=O of PyEG, enhancing the interfacial interactions inducing the transformation from α-­phase to β-­phase PVDF crystal. The synergistic effect of conjugation degree and interfacial interactions causes a remarkable increase in dielectric constant of 480 at 1 kHz, 55 times higher than that of pure PVDF for the presence of 0.74 vol% filler in the composite. This study provides a novel perspective to tailor polymer composites with high dielectric constant and low loss by regulating the conjugation degree of graphene-­based nanofillers and their interfacial interactions with polymer matrix. Long et al.51 used the noncovalent functionalization of graphene by solution-­phase assembly of aminodecane, where the alkane chains remain oriented perpendicular to the graphene basal plane. Raman spectra suggest that the self-­assembly occurs by a noncovalent process, i.e. the sp2 hybridization of graphene remains intact. The self-­assembly of 1-­aminodecane on graphene is a potential route for passivating graphene field-­effect devices. Polyoxometalates (POMs) are the molecular clusters of a transition metal oxide, displaying a large number of structural varieties having use in catalysis, medicine, magnetics, electronics and also in fabricating hybrid polymers.52

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Wang et al. used a noncovalent phase-­transfer-­assisted method to fabricate polymer functionalized graphene.53 They synthesized a series of cluster-­cored star polymers (CSPs) containing polyoxometalate core and polystyrene (PS) arms which act as modifiers. A Keggin-­t ype POM, H4SiW12O40 (SiW12), was selected as the inorganic core, and polystyrene (PS) arms are then grafted by RAFT polymerization, from the chain transfer agent (CTA) premodified on SiW12 (Figure 3.13). By adjusting the molar ratios of styrene to SiW12 (100 : 1, 200 : 1, to 300 : 1) three CSPs, CSP-­1, CSP-­2 and CSP-­3, respectively are synthesized with the same SiW12 core but with different chain length of PS arms. The CSPs can strongly adsorb on graphene nanosheets by the electron transfer interaction between graphene nanosheets and polyoxometalate. The CSPs undergoes phase transfer from aqueous to organic media like chloroform, producing individually dispersed graphene. The CSP-­functionalized graphene is compatible and well dispersible with additional PS matrices and serves as a reinforcing nanofiller for polymer composites. It has been found that a 0.2 wt% loading of CSP-­3 in a PS matrix achieves a 98.9% high increase in the Young’s modulus.

Figure 3.13  Phase-­  transfer-­assisted fabrication of CSP functionalized RGO and its

PS composite coating. Reproduced from ref. 53 with permission from American Chemical Society, Copyright 2015.

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To achieve high mechanical and electrical properties in epoxy resin Li et al.54 used mixed π–π stacking and H-­bonding interactions to make rGO/epoxy composite using polystyrene sulphonate (PSS) as exfoliating agent. Here GO is chemically reduced with hydrazine hydrate in the presence of PSS in aqueous medium to produce a rGO/PSS hybrid via π-­stacking interaction. The –SO3H group of PSS helps the exfoliated rGO sheets to make good dispersion of the PSS-­rGO hybrid, facilitating nanocomposite preparation with epoxy polymer. The composite is made by adding epoxy resin to the PSS-­rGO dispersion in alcohol and the mixture is ultrasonicated and finally evaporated. The noncovalent functionalization increases the interfacial bonding between the epoxy matrix and graphene via H-­bonding through the hydrogen atoms of unreduced oxygenated groups of GO with the epoxide groups of the polymer resin. This strong interfacial bonding in the composite for 1 wt% filler leads to enhance tensile modulus to 2525 MPa, which is 28% higher, and tensile strength reaches 89.8 MPa, 21.8% higher than that of neat epoxy. The PSS-­rGO addition enhances dc-­conductivity by six orders for 0.73 vol% addition of the filler and the dielectric constant of the epoxy/2 wt% PSS-­g nanocomposite at 50 Hz is about 881, which is much higher than that of neat epoxy.

3.6  M  ixed Noncovalent and Covalent Functionalization In order to achieve high performance graphene/polymer composite sometimes a combination of noncovalent and covalent functionalization is used by different groups of workers. Chen et al.55 prepared a high performance graphene-­based polyurethane composite by exploiting covalently and noncovalently pyrene functionalized rGO. The residual hydroxyl and epoxide groups of rGO help to make a covalent bond with polyurethane chains by sequentially reacting with diisocyanate and polyethylene glycol oligomer. The noncovalently bonded polyurethane chains are produced by the π–π interaction between rGO and pyrene derivatives. This mixed functionalized structure produces a tough polymer composite with high ductility for the presence of H-­bonds between the polar functional groups and π–π interaction between aromatic elements. The former interaction improves the dispersion of graphene in the polyurethane matrix and the later enhances the efficiency of interface load transfer. In a similar way by exploiting combined noncovalent and covalent bonds, a highly efficient anti-­corrosive epoxy coating, comprising homogeneously dispersed graphene (Gr), was achieved by He et al.56 using tannic acid (TA) as intercalator to disperse Gr in water. TA was chosen for its phenyl and hydroxyl groups, which not only afford strong π–π noncovalent interaction between Gr and TA but also facilitate good dispersion of Gr–TA hybrids in aqueous medium. So, γ-­(2,3-­epoxypropoxy) propyltrimethoxysilane (KH560) is used to modify the surface of Gr–TA in which the –OH groups of TA react with the Si–OH groups of KH560 being produced on hydrolysis. Simultaneously, the epoxy groups of KH560 react with the epoxy groups of the epoxy

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via a covalent bond to enhance the interactions between Gr–TA and the epoxy (Figure 3.14). Aqueous epoxy resin is fully mixed with the Gr–TAKH560 dispersion, and is dried a via rotary evaporator. Finally the system is cured with a curing agent and is painted onto the pretreated metal substrates by high pressure

Figure 3.14  The  reaction progress of Gr-­TA-­KH560 hybrid material. Reproduced from ref. 56 with permission from American Chemical Society, Copyright 2019.

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spraying. The corrosion resistance properties of the composite Gr–TA/epoxy coatings are tested electrochemically and the results indicate that the anti-­ corrosion properties of Gr–TA–KH560/epoxy composite coating is highly enhanced from those of the Gr–TA/epoxy composite and pure epoxy polymer.

3.7  Scope Thus noncovalent functionalization of graphene with polymers can be achieved using varieties of interaction like π–π, H-­bonding, ionic, hydrophobic and electron transfer interactions. This nonbonding interaction of graphene is superior to covalent interactions because the conjugation of graphene is not at all hampered and this enables good conductivity and optoelectronic properties. Futher, the higher aspect ratio of graphene in noncovalent functionalization, compared to the covalent functionalization, also yields high mechanical properties in the composite. The important feature of noncovalent functionalization is that it can be further functionalized with covalent functionalization to produce a technologically important polymer composite. Another important scope of this functionalization technique is that one can adopt a procedure to immobilize metal and semiconducting nanoparticles to impart new optoelectronic properties giving an immense scope for further work for material scientists and engineers to develop new exciting materials suitable for different applications, including biotechnological applications.

References 1. Y. X. Xu, H. Bai, G. W. Lu, C. Li and G. Q. Shi, J. Am. Chem. Soc., 2008, 130, 5856–5857. 2. Q. H. Wang and M. C. Hersam, Nat. Chem., 2009, 1, 206. 3. Q. Su, S. Pang, V. Alijani, C. Li, X. Feng and K. Mullen, Adv. Mater., 2009, 21, 3191. 4. J. R. Lomeda, C. D. Doyle, D. V. Kosynkin, W. F. Hwang and J. M. Tour, J. Am. Chem. Soc., 2008, 130, 16201. 5. J. Liu, W. Yang, L. Tao, D. Li, C. Boyer and T. P. Davis, J. Polym. Sci., Part A: Polym. Chem., 2010, 48, 425. 6. I. Liu, L. Tao, W. Yang, D. Li, C. Boyer, R. Wuhrer, F. Braet and T. P. Davis, Langmuir, 2010, 26, 10068. 7. H. Li, L. Han, J. J. Cooper-­White and I. Kim, Nanoscale, 2012, 4, 1355. 8. L. Xu, L. Wang, B. Zhang, C. H. Lim, Y. Chen, K. Neoh, E. Kang and G. D. Fu, Polymer, 2011, 52, 2376. 9. K. Jo, T. Lee, H. J. Choi, J. H. Park, D. J. Lee, D. W. Lee and B. Kim, Langmuir, 2011, 27, 2014. 10. H. Bai, Y. Xu, L. Zhao, C. Li and G. Shi, Chem. Commun., 2009, 13, 1667. 11. X. Qi, K. Pu, H. Li, X. Zhou, S. Wu, Q. Fan, B. Liu, F. Boey, W. Huang and H. Zhang, Angew. Chem., Int. Ed., 2010, 49, 9426.

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12. L. Peponi, A. Tercjak, R. Verdejo, M. A. Lopez-­Manchado, I. Mondragon and J. M. Kenny, J. Phys. Chem. C, 2009, 113, 17973. 13. S. Stankovich, R. D. Piner, X. Chen, N. Wu, S. T. Nguyen and R. S. Ruoff, J. Mater. Chem., 2006, 16, 155–158. 14. L. Ji, Y. Wu, L. Ma and X. Yang, Composites, Part A, 2015, 72, 32–39. 15. C. Chen, S. Qiu, M. Cui, S. Qin, G. Yan, H. Zhao, L. Wang and Q. Xue, Carbon, 2017, 114, 356. 16. M. Lian, J. Fan, Z. Shi, H. Li and J. Yin, Polymer, 2014, 55, 2578. 17. N. Maity, A. Mandal and A. K. Nandi, Polymer, 2016, 88, 79. 18. J. Kim, J. Cha, B. Chung, S. Ryu and S. H. Hong, Compos. Sci. Technol., 2020, 192, 108101. 19. J. Cho, I. Jeon, S. Y. Kim, S. Lim and J. Y. Jho, ACS Appl. Mater. Interfaces, 2017, 9, 27984. 20. S. Song, C. Wan and Y. Zhang, RSC Adv., 2015, 5, 79947. 21. W. Chen, P. Liu, L. Min, Y. Zhou, Y. Liu, Q. Wang and W. Duan, Nano-­Micro Lett., 2018, 10, 39. 22. H. Hu, Z. Hu, X. Ren, Y. Yang, R. Qiang, N. An and H. Wu, Chin. J. Chem., 2015, 33, 199. 23. L. Hou, J. Gao, H. Ruan, X. Xu and S. Lu, J. Polym. Res., 2020, 27, 32. 24. F. Liu, J. Y. Choi and T. S. Seo, Chem. Commun., 2010, 46, 2844. 25. F. Sharifi, R. Bauld, M. S. Ahmed and G. Fanchini, Small, 2012, 8, 699. 26. Y. Liu, R. Li, B. Liang, C. Li, J.-­H. Hu, K. Zeng and G. Yang, J. Mater. Sci., 2020, 55, 140. 27. J. Liang, Y. Huang, L. Zhang, Y. Wang, Y. Ma, T. Guo and Y. Chen, Adv. Funct. Mater., 2009, 19, 2297. 28. R. K. Layek, S. Samanta and A. K. Nandi, Carbon, 2012, 50, 815. 29. R. K. Layek, S. Samanta and A. K. Nandi, Polymer, 2012, 53, 2265. 30. R. K. Layek, A. Kundu and A. K. Nandi, Macromol. Mater. Eng., 2012, 298, 1166. 31. X. Wang, H. Bai, Z. Yao, A. Liu and G. Shi, J. Mater. Chem., 2010, 20, 9032. 32. S. Zu and B. Han, J. Phys. Chem. C, 2009, 113, 13651. 33. E. Choi, T. H. Han, J. Hong, J. E. Kim, S. H. Lee, H. W. Kim and S. O. Kim, J. Mater. Chem., 2010, 20, 1907. 34. S. Liu, J. Tian, L. Wang, H. Li, Y. Zhang and X. Sun, Macromolecules, 2010, 43, 10078. 35. A. Kundu, R. K. Layek and A. K. Nandi, J. Mater. Chem., 2012, 22, 8139. 36. A. Kundu, R. K. Layek, A. Kuila and A. K. Nandi, ACS Appl. Mater. Interfaces, 2012, 4, 5576. 37. A. M. Pandele, C. Andronescu, E. Vasile, I. C. Radu, P. Stanescu and H. Iovu, Composites, Part A, 2017, 103, 188–195. 38. B. Gupta, N. Kumar, K. Panda, A. A. Melvin, S. Joshi, S. Dash and A. K. Tyagi, J. Phys. Chem. C, 2016, 120, 2139. 39. R. K. Layek, E. Uddin, N. H. Kim, A. Kin, T. Lau and J. H Lee, Composites, Part B, 2017, 128, 155. 40. S. Park, J. An, J. R. Potts, A. Velamakanni, S. Murali and R. S. Ruoff, Carbon, 2011, 49, 3019–3023.

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Chapter 4

Physical Properties of Polymer Functionalized Graphene 4.1  Morphology The addition of nanoparticles into the polymer matrix influences the packing of polymer chains, and this is highly affected if there is interaction between the components. Graphene, functionalized graphene and GO can yield different interactions with polymers, changing the organization and packing of chains, and hence effect the overall texture of the PFG. We shall discuss it by studying it using transmission, scanning electron and atomic force microscopic images in the following sections.

4.1.1  Transmission and Scanning Electron Microscopy Transmission electron microscopy (TEM) is an essential technique to characterize PFG. Fang et al.1 covalently attached initiator molecules on the graphene surface by diazonium addition followed by ATRP polymerization of styrene chains to the graphene nanosheets. They observed a density difference between the central and peripheral regions of polystyrene (PS) functionalized graphene sheets and found the average lateral size of PS functionalized graphene is about 100 nm. A colour difference between the central and peripheral areas of each PFG sheets was also noticed by them. Yang et al.2 grafted poly(dimethyl amino ethyl methacrylate) (PDMAEMA) from the GO surface by the ATRP technique and TEM study revealed that polymer particles are decorated on the GO sheets as dark nanodots (Figure 4.1).

  Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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Figure 4.1  TEM  image of PS-­grafted graphene sheets. Reproduced from ref. 1 with permission from the Royal Society of Chemistry.

Figure 4.2  TEM  images of (a) pristine GO and (b) GO-­g-­P(PEGMA). Reproduced from ref. 4 with permission from Elsevier, Copyright 2017.

On the other hand, Li et al.3 synthesized poly(tert-­butyl acrylate) (PtBA) brushes by surface-­initiated ATRP, and the sheet morphology of GO remained unchanged in the PtBA functionalized GO. Wang et al.4 observed from TEM study that when GO is covalently functionalized with poly(ethylene glycol) methyl ether methacrylate (PEGMA) it overcomes the aggregation of GO nanosheets (Figure 4.2a and b). For the pristine GO nanosheets, quite a few layers of assemble together to form a semi-­transparent multilayer aggregate. In contrast, from Figure 4.2b, it appears that GO-­g-­P(PEGMA) is in the form of a single layer with a rougher and more opaque state. It indicates that the agglomeration of the GO nanofillers has been alleviated. The particle size distribution of GO and GO-­g-­P(PEMGA) nanosheets decreased from 1246 nm for GO to 428 nm for GO-­g-­P(PEGMA) after the modification. Also the particle size distribution of GO-­g-­P(PEMGA) becomes narrower in range, suggesting that the large particle of agglomerated GO becomes diminished after grafting of P(PEGMA). The morphology of 2-­hydroxyethyl methacrylate (HEMA) grafted GO using

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ATRP modified graphene samples, as observed by TEM study, is presented in Figure 4.3.5 The initiator attached to GO (Init-­G) exhibits flake-­like morphology with wrinkles and pleats. After polymer functionalization, PHEMA-­G becomes well coated with PHEMA (Figure 4.3b), as evident from the ‘dark clouds’ of the amorphous PHEMA chains on the graphene surface.5 Interestingly, the PHEMA-­G has greater polymer density at its edges, which it may be due to the amorphous edge carbon atoms of GO being more easily attacked and modified by PHEMA. Santos et al.6 esterified carboxylic groups of nanoGO with the glycol containing porphyrin and analysed morphologies with different chain length of glycolic unit by TEM (Figure 4.4). The TEM images of these materials in Figure 4.4B and C show darker areas in the hybrids in a higher amount compared with pristine nanoGO sheets (Figure 4.4A). Their surface exhibits a wrinkled structure having scrolled edges and broad dark spots. These dark spots are attributed to some degree of sheet folding resulting from coupling of nanoGO with the glycol porphyrins and additionally to aggregation induced by the porphyrin brushes. This indicates the successful attachment of the glycol porphyrin molecules on nanoGO-­CO2H sheets.

Figure 4.3  Representative  (a) TEM image of Init-­G, (b) TEM image of PHEMA-­G. Reproduced from ref. 5 with permission from American Chemical Society, Copyright 2017.

Figure 4.4  TEM  images of nanoGO (A), nanoGO-­P2 (B) and nanoGO-­P5 (C). Reproduced from ref. 6 with permission from Elsevier, Copyright 2018.

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Turlakov et al. functionalized GO with a series of conjugated phenyleneethynylene (PPE) copolymers bearing different electron-­donating and/ or withdrawing groups. The HRTEM images of GO-­PPE hybrids (Figure 4.5) show that the molecules self-­assemble in blocks resembling bricks like shapes that can be described as sanidic liquid crystal (LC) materials to indicate their mesomorphism. Functionalization is mainly carried out at the edges of the GO sheets, so that the GO sheets are not totally covered, because of the tendency of the PPEs to self-­assemble in an ‘edge-­on’ conformation, indicating that the conjugated backbones are parallel to the GO surface, rather than ‘face-­on’, where the backbones are flat-­lying on the GO surface. The calculated periodicity 0.33 nm coincided with the distance between conjugated benzoates. They observed that the molecules are packed in random blocks, as confirmed by their selected area electron diffraction patterns where many diffraction spots with the same periodicity are seen. Xie et al. 8 grafted polyacrylate introducing an intermediate l3-­ methacryloxypropyl trimethoxysilane (KH-­570) onto the GO surface followed by polymerization of acrylate monomer (60 wt% methacrylate and 40 wt% glycidyl methacrylate) through in situ free radical random polymerization on the surface of GO sheets. The original GO shows a highly transparent and smooth nanosheet structure (Figure 4.6a) and after modification with KH-­570 silane, the GO sheets exhibited significant stretching, indicating the incorporation of KH-­570 reduced the layer–layer interaction, producing lot of wrinkles (Figure 4.6b). After grafting of polyacrylate copolymer, the PA-­GO sheets become less transparent, showing dark features (Figure 4.6c), which indicates the surface of the GO sheets becomes surrounded by a thin polyacrylate layer and the layer structure of GO sheets has not been changed in PA-­GO after covalent modification.

Figure 4.5  (a)  HRTEM image of a casting film of EC copolymer deposited on a

lacy carbon grid showing a periodicity of 0.32 nm and corresponding to the distances between conjugated backbones, (b) HRTEM image of the same film and (c) sketch of the molecular arrangement seen from a lateral view. Reproduced from ref. 7 with permission from Elsevier, Copyright 2017.

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Figure 4.6  TEM  images of GO (d), f-­GO (e) and PA-­GO (f). Reproduced from ref. 8 with permission from Elsevier, Copyright 2020.

Figure 4.7  HR-­  TEM images of (a) rGO and (b) PFG. Reproduced from ref. 9 with permission from Elsevier, Copyright 2016.

The morphology of noncovalently functionalized polythiophene-­g-­ polymethyl methacrylate (PT-­g-­PMMA) with rGO, produced by π-­stacking interaction, is monitored from TEM images (Figure 4.7a and b). From the figure it is observed that rGO sheets are in a wrinkled state and in the PFG the light dark phase on the edge and on the surface of the rGO of wrinkled sheets is attributed to the noncovalent attachment of a PT-­g-­PMMA chain.9 The resulting PFG forms stable dispersion even after more than one month under ambient conditions. The TEM micrographs of sulfonated graphene (SG) (Figure 4.8) indicate the sheet-­like morphology of SG and the dendritic morphology of SG3 (the number denotes the weight percentage of SG in its composite with PVA).10 In SG3 the sheet morphology of SG is totally absent, rather the fibrils forms dendrites, which is better clarified in the FESEM micrographs (Figure 4.9) where SG3 shows dendritic and SG5 shows rod-­like morphology, but none of these morphologies are present in pure SG and PVA. The new morphology may arise from the self-­organization of the supramolecular complex between SG and PVA producing fibrils and hence dendrites. The evidence of supramolecular interaction (H-­bonding) is evident from FTIR study. The amorphous layer of SG–PVA together with PVA lamella forms a supramolecular complex which organizes to produce fibrils that grow longitudinally and laterally producing different morphology depending on the composition. In SG3 the

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Figure 4.8  TEM  images of (a) SG and (b) SG3 samples. Reproduced from ref. 10 with permission from Elsevier, Copyright 2011.

Figure 4.9  FESEM  micrographs of (a) SG3 and (b) SG5 composites. Reproduced from ref. 10 with permission from Elsevier, Copyright 2011.

interfacial interactions are conducive to the formation of special dendritic morphology.10 Li et al.11 studied the morphology of perylene tetracarboxylic acid (Py) anchored exfoliated graphene (PyEG) nanosheets within the PVDF matrix using FESEM (Figure 4.10). Here, the freeze-­fractured microstructures of EG/ PVDF and PyEG/PVDF composites exhibit slight aggregation of EG in the PVDF matrix (Figure 4.10a). However, PyEG/PVDF exhibits homogeneous dispersion of PyEG without any aggregation, showing better compatibility between PyEGnanosheets and the PVDF matrix due to the enhanced interfacial interactions through the –COOH groups (Figure 4.10b). This enhanced distribution of PyEG nanosheets is due to the enhancement of interfacial interactions with the PVDF matrix facilitating the formation of microcapacitors in the hybrid. The elements mapping images of Figure 4.10c for PyEG(0.5%)/PVDF exhibit uniform distributions of carbon, fluorine and oxygen, supporting the

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Figure 4.10  Cross-­  section SEM images of (a) EG(0.5 vol%)/PVDF, (b) PyEG(0.5 vol%)/PVDF composites and (c) corresponding elements mapping images of C, F and O in the area. Reproduced from ref. 11 with permission from Elsevier, Copyright 2018.

Figure 4.11  Scanning  electron microscopy (SEM) images of the tensile failure section for PLLA (a) GO-­g-­PLLA3/PLLA (b) and GO-­g-­PLLA5/PLLA (c). Image in the red circle demonstrates the aggregation of nanofiller. Reproduced from ref. 12, https://doi.org/10.3390/polym9090429, under the terms of the CC BY 4.0 license, https://creativecommons. org/licenses/by/4.0/.

homogeneous distribution of PyEG nanosheets in the PVDF matrix by noncovalent interaction. GO-­g-­poly(l-­lactic acid) (GO-­g-­PLLA) produced by condensation reaction at 120 °C is used to make GO-­g-­PLLA/PLLA nanocomposites using a solvent mixing method. The SEM images of fracture surfaces of PLLA, GO-­g-­PLLA3, and GO-­g-­PLLA5 (Figure 4.11) show that the fracture surface of PLLA is much smoother than that of composites due to the interfacial interaction between the filler and matrix PLLA. This causes an increase in tensile strength and

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the tensile fracture surface become coarser, showing some fibres on the fracture surfaces of the GO-­g-­PLLA3 nanocomposite. This arises due to stress transfer from PLLA to GO sheets, suggesting good interaction and compatibility between GO-­g-­PLLA and the PLLA matrix. It is evident from Figure 4.11, that the fractured surface of Go-­g-­PLLA5/PLLA is much coarser and no fibril structure has appeared; instead some aggregates are observed indicating that GO-­g-­PLLA has not dispersed well, rather it has become aggregated at higher filler concentration.

4.1.2  Atomic Force Microscopy (AFM) AFM image and the corresponding height profile is an important tool to characterize the polymer functionalization of graphene sheets. The AFM study offers information about the length and thickness of GO along with the morphology of the PFG. For AFM study the GO sample in water (0.01%) is cast on a freshly cleaved mica surface and is observed through the instrument. The sheet thickness is measured from AFM height profile and similarly for the initiator or polymer functionalized samples.1,13 The height of the initiator and polymer functionalized GO increases from that of pristine GO. Two major populations of height are observed in PMMA functionalized GO (MG), i.e. 2 and 4.5 nm (Figure 4.12) and this happens possibly due to polymerization at the side and basal planes of GO.13 The 2 nm height may arise for the polymerization at the initiator on the side of GO and PMMA thus coating the graphene sheet and increasing the thickness. The occurrence of 4.5 nm thickness may be due to polymerization from the initiator at the basal plane of the GO sheet. PMMA, thus produced, enters into the gallery of graphene sheets and increases the gallery spacing to yield 4.5 nm thickness. In the MG/ PVDF composite (MG3)13 the specific interaction between PVDF and PMMA chains facilitates entry of PVDF chains into the MG gallery, expanding the MG gallery by reduction of the cohesive force between the two graphene sheets. As a result, exfoliation of MG sheets occurs showing lower height of MG3 in the AFM height images. It may be inferred that 2 nm height populated MG breaks into 0.5, 1 and 1.5 nm, and the 4.5 nm height population exfoliates into 3 nm and lower height population. Thus different heights of MG sheet occur due to exfoliation at different levels of the graphene sheets during blending of MG with PVDF. Ding et al.14 synthesized polystyrene from the surface of thermally reduced graphene oxide (TRGO) grafted reversible addition–fragmentation chain transfer (RAFT) agent (TRGO-­RAFT) as a ‘core’ material. The AFM results of TRGO-­RAFT and PRG composites are shown in Figure 4.13. The RAFT agents are grafted onto the basal planes of graphene via esterification reaction and the thickness of the TRGO-­RAFT sheets increase from 0.8 nm (TRGO) to 2.6 nm (Figure 4.13a), evidencing successful surface RAFT agent grafting. The raft polymerization of styrene onto the basal planes of graphene causes an increase in the height of the PRG sheets to 4.9 nm (Figure 4.13b), evidencing the successful surface grafting of polymer. It indicates that the TRGO-­RAFT

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Figure 4.12  Tapping  mode AFM images and height profile of BIBG (a), MG (b)

and MG3 (c) and corresponding histogram on the right-­hand side (obtained from ∼30 objects). Reproduced from ref. 13 with permission from Elsevier, Copyright 2010.

precursor acts as a ‘core’ material to prepare PS/graphene composites. On comparing the AFM images of TRGO-­RAFT and PRG composites, it is found that the size of graphene became smaller when polymer chains grow on the surface of graphene. In noncovalent polymer functionalized graphene similar increase of thickness of graphene is observed. Ji et al.15 studied the AFM morphology of pyrene-­terminated liquid crystalline polymer (LCP) functionalized graphene produced via π–π interactions (Figure 4.14). The greater surface roughness of LCP functionalized GO compared to GO is evident from the AFM images. The height profile of GO indicates the average thickness of GO sheets is ∼0.8 nm, similar to the previous results.16,17 As the LCP is attached to the graphene

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Figure 4.13  AFM  images of (a) TRGO-­RAFT and (b) PRG. Reproduced from ref. 14 with permission from Elsevier, Copyright 2014.

surface via π–π interactions, the polymer covers both sides of graphene sheet increasing the thickness to 2.2 nm. Maity et al.9 studied the topological AFM images of rGO and PFG (Figure 4.15) and it is evident from the height profile of GO that GO has thickness of 0.84 nm, characterizing almost a single layer of GO. On noncovalent functionalization with PT-­g-­PMMA the height profile of PFG indicates two different heights (1.41 and 2.21 nm), which are higher than that of GO. The former height can be attributed to the monolayer polymer functionalization on the rGO sheet and the latter may arise from noncovalent grafting surrounding both the surfaces of rGO sheet. Thus, the AFM height images confirm noncovalent functionalization of rGO with PT-­g-­PMMA. Layek et al.18 made a hybrid of sulphonated graphene and chitosan for noncovalent functionalization. Chitosan is soluble in water/acetic acid medium and possesses many hydroxyl and amino groups accessible for interaction with SG facilitating molecular level dispersion of SG in the chitosan matrix due to H-­bonding interaction between the NH2 group and –SO3H, –COOH and –OH groups of SG. In Figure 4.16 a tapping-­mode AFM image of SG sheets and the height profile analysis from a large number of SG images exhibits that the average height of the SG sheets is 2.04 nm and average length is 1015 nm. This indicates that SG sheets, dispersed as a few (∼2–3) exfoliated graphene sheets, are attached together in the chitosan matrix. From the comparison of AFM images of the pure chitosan and SG5 samples the SG sheets (white patches) are clearly observed, indicating that they are dispersed homogeneously in the grey chitosan matrix. They also made hybrids of SG and PVA via H-­bonding interaction and found that the increase in height in the PVA/SG composite depends on the concentration of filler forming different texture.10

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Figure 4.14  AFM  image and cross-­section contour of (A) GO sheets; (B) LCP (BP6-­ 6400 g mol−1) functionalized graphene sheets. Reproduced from ref. 15 with permission from Elsevier, Copyright 2015.

4.2  Structural Study The structural study of PFGs is usually made from FTIR, Raman and XPS spectroscopy as well as with wide angle X-­ray scattering and are discussed both for covalent and noncovalent polymer functionalization of graphene.

4.2.1  Fourier Transformed Infrared Spectroscopy (FTIR) FTIR spectra is an important tool to understand both the covalent and noncovalent polymer functionalization of GO. The pristine graphite has a characteristic band at 3430 cm−1 for O–H stretching and at 1610 cm−1 for skeletal vibrations from graphitic domains of adsorbed water and aromatic domain (C=C), respectively (Figure 4.17).19 The GO shows some new peaks at 1785/1707 cm−1

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Figure 4.15  Tapping  mode AFM images of (a) GO and (b) PFG with their height

profiles. Reproduced from ref. 9 with permission from Elsevier, Copyright 2016.

Figure 4.16  (a)  Tapping-­mode AFM image of SG and (b) height profile (c) length

profile (d) histogram of height profile and (e) histogram of length profile of SG (obtained from w20 objects). Reproduced from ref. 18 with permission from Elsevier, Copyright 2012.

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Figure 4.17  FTIR  spectra of graphite, GO and RGO. (a) GO; (b) RGO (95 °C for 3 h) and (c) graphite. Reproduced from ref. 19 with permission from Elsevier, Copyright 2017.

(C=O stretching vibrations from carbonyl and carboxylic groups), 3430 cm−1 for O–H stretching, and 1105 cm−1 for C–O stretching and at 2925/2855 cm−1 (CH/CH2) assigned to methylene stretching representing the existence of some CH2 or CH groups in the GO samples. The C=C aromatic peak shows a characteristic peak at 1610 cm−1 in graphite and it is shifted to 1625 cm−1 in GO due to the presence of some electron-­withdrawing oxygen-­containing functional groups. The rGO shows a characteristic peak of a hydroxyl group (3430 cm−1) with reduced intensity compare to that of GO but the characteristic band of epoxide group (1610 cm−1) becomes totally disappeared.19 Lin et al.20 covalently bonded polyethylene grafted GO by firstly coating gamma-­ aminopropyl triethoxysilane (APTES) onto the GO sheets, followed by grafting with maleic anhydride grafted polyethylene (MA-­g-­PE) and characterized by means of FTIR spectroscopy. The presence of FTIR absorption peaks at 1045 and 1119 cm−1 for the Si–O–C and Si–O–Si vibrations indicate successful coating of the APTES onto GO through chemical bonding. The occurrence of peak at 1547 and 802 cm−1 in polyethylene grafted GO represent secondary amide N–H bending and C–N stretching, indicating the presence of the primary amide bonding (–CO–NH–). These results provide evidence of covalent grafting of the polyethylene onto the APTES coated GO through the chemical reaction of the –NH2 group of APTES with the maleic anhydride of MA-­g-­PE. Wang et al.21 used the SI-­ATRP technique to graft poly(ethylene glycol) methyl ether methacrylate (PEGMA) on GO. The prepared GO, GO-­NH2, GO-­Br and GO-­g-­P(PEGMA) nanoplates were characterized by FTIR spectra (Figure 4.18), and the pristine GO exhibits three FTIR peaks at 1050, 1700 and 3420 cm−1, indicating the presence of C–O, COO− and OH groups on the GO surface, respectively. The disappearance of the –OH vibration peak at 3420 cm−1 and the presence of two new peaks at 1080 cm−1 (Si–O)and 2930 cm−1 (C–H) suggest successful grafting of aminopropyltriethoxysilane (APTES) on the GO surface

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Figure 4.18  FTIR  spectra of GO, GO–NH2, GO–Br and GO-­g-­P(PEGMA). Reproduced from ref. 21 with permission from Elsevier, Copyright 2017.

producing GO-­NH2. The two sharp peaks at 1630 and 1530 cm−1 of the N–C=O stretching vibration suggest the successful reaction between GO-­NH2 and 2-­bromoisobutyryl bromide, producing the ATRP initiator. On ATRP polymerization of PEGMA the formation of GO-­g-­P(PEGMA) is evident from the enhanced absorption peaks at 2870 and 1700 cm−1 for the CH3 and carboxylic groups, respectively, for the absorption peaks of P(PEGMA) suggesting that P(PEGMA) is successfully grafted on the GO surface. In the case of noncovalent functionalization of GO with poly(vinyl alcohol) (PVA) Layek et al. found that the –OH stretching peak at 3424 cm−1 of GO shifts to lower energy at 3265 cm−1 for the GO-­PVA hybrid (2 wt% GO) suggesting H-­bonding interactions between PVA and GO sheets.5 The 1710 cm−1 peak of >C=O group shifts to 1727 cm−1 and the epoxide stretching vibration of GO at 1060 cm−1 shifts to 1100 cm−1 indicating H-­bonding between the –OH group of PVA and the epoxy group of GO. The increased frequency of vibration is due to the ring structure of the H-­bonded epoxy group of GO in the GO-­PVA hybrids. Maity et al.22 successfully oxidized graphite to GO showing a broad peak of the O–H group (Figure 4.19a) at 3425 cm−1, strong C=O stretching peak at 1721 cm−1, in-­plane deformation vibration of the O–H group at 1388 cm−1, C–OH stretching peak at 1242 cm−1, epoxide C–O stretching peak at 1058 cm−1 and the 1624 cm−1 peak is for the unoxidized graphitic frame work. Reduction of GO with hydrazine leads to the removal of the >C=O peak suggesting the absence of oxygen functionalities including the carboxylic acid group. The occurrence of the 1634 cm−1 (C=C) peak indicates the restoration of the π-­conjugated network of graphene and the shift of the graphitic vibration to higher energy indicates the formation of many small size graphitic domains during the reduction. In PT-­g-­PMMA the 1733 cm−1 peak is due to the >C=O group of PMAA and in the PT-­g-­PMMA functionalized rGO the C=C vibration of the graphene skeleton shifts to 1627 cm−1 from that of rGO (1634 cm−1) for the noncovalent functionalization. This shift to lower energy

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Figure 4.19  (a)  FTIR spectra of GO, RGO, PT-­g-­PMMA, PFG and (b) Raman spectra

of graphite, GO, RGO, PFG in a solid sample. Reproduced from ref. 22 with permission from Elsevier, Copyright 2016.

indicates that the π-­electrons of the graphitic skeleton become more delocalized due to π–π interaction of rGO with the thiophene units in PT chains of PT-­g-­PMMA in the PFG sample.

4.2.2  Raman Spectroscopy All the carbonaceous materials, e.g., graphene, carbon nanotubes, nanodiamonds, etc. display high Raman intensities, suggesting them as a potential tool for characterization of functionalized graphenes.23 The D-­band of GO occurs at 1366 cm−1 and is attributed to the stretching of sp3 carbon of the graphene sheets and the stretching of sp2 carbon occurs at 1582 cm−1 forming the G-­band.24,25 The intensity ratio of the D and G bands (ID/IG) yields a measure of the extent of disorder present within the graphene26 and the ID/IG ratio changes on functionalization of graphene. During covalent functionalization of GO to produce molecularly imprinted GO/polymethacrylamide (PMAAM) hybrid (GO-­MIP) material by the ATRP technique, Chang et al.27 observed ID/IG values of 0.670, 0.884 and 1.169 for GO, GO–Br, and GO–MIP, respectively. These results indicate an increase in disorder due to modification of graphene and as a long polymer chain is introduced the ID/IG increases more. Not only the ID/IG ratio, but also the Raman shift of D and G band positions can give information about the interaction between polymer and GO,28 and the shift of the G-­band is attributed to the alteration of the electronic structure of graphene. Usually the electron donor dopants shift the G-­band to lower frequencies and electron acceptor dopants shifts to the higher frequencies.29,30 This is due to the fact that the electron donor dopants cause increased electron density in the graphene, facilitating better delocalization of its π-­electrons shifting the G-­band to lower frequency. However, for electron acceptor dopants, the opposite behaviour occurs as the ring electrons become localized, making the stretching vibration of sp2

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carbon of graphene difficult. Fang et al. covalently functionalized the GO surface via diazonium addition followed by ATRP to link polystyrene chains (82 wt%) on the GO nanosheets. From Raman spectra the ID/IG ratio of GO, initiator-­grafted GO, and PS-­grafted GO sheets are found to be 1.02, 1.71 and 1.73, respectively (Figure 4.20), reflecting an increase in the disorder after diazonium addition. Such disorder is also reflected in the broadened and blue-­shifted G bands for the GO appearing at 1595 cm−1, which is 15 cm−1 higher than that of pristine graphite (1580 cm−1). After reduction of GO with hydrazine, the vibration frequency of the G band decreases to 1585 cm−1, with an increased ID/IG ratio relative to GO.31 The change in the D/G ratio can be partly attributed to the formation of covalent bonds between graphene and the initiator and/or PS chains. Apart from the covalent functionalization, in noncovalent polymer functionalization a similar change in the Raman spectra is also noticed. Kundu et al.32 functionalized poly(vinyl alcohol) (PVA) with GO via H-­bonding interaction and thoroughly characterized it by Raman spectroscopy by making the samples at two different pHs. The Raman spectrum of GO (Figure 4.21a) at pH 4 exhibits the D band (doublet) at 1388 and 1469 cm−1 and the G band occurs at 1613 cm−1 with ID/IG ratio 1.01. In the GO-­PVA hybrids at pH 4, both D and G bands shift to a lower wave number and the ID/IG ratio increases to 1.14 and 1.28 for GO-­PVA0.5 and GO-­PVA1, respectively, with increasing PVA concentration. Both the D and G bands arise due to skeletal vibrations of GO at the sp3-­and sp2-­rich domains, respectively. After complexation with PVA, the π-­electrons of graphene rings become somewhat delocalized, facilitating vibration of the skeleton and shifting the above bands to lower energy due to increased electron density arising from the nonbonding electrons of the

Figure 4.20  Raman  spectra of the pristine graphite (1), graphite oxide (2), graphene-­ initiator (3) and graphene-­PS (4), with a laser excitation wavelength of 631 nm. Reproduced from ref. 1 with permission from the Royal Society of Chemistry.

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Figure 4.21  Raman  spectra of GO, GO-­PVA0.5, and GO-­PVA1 (a) at pH 4 and (b) at

pH7. Reproduced from ref. 32 with permission from American Chemical Society, Copyright 2012.

–OH group of PVA. Further an increase of ID/IG ratio with PVA concentration indicates increased sp3 character because of increased hydrogen bonding of GO with PVA. To elucidate the origin of the doublet of the D band of GO, GO-­PVA0.5, GO-­PVA1 at pH 4 it is necessary to compare with those at pH 7 (Figure 4.21b), where the D band exhibits a singlet in the range 1352–1355 cm−1. This lower energy skeleton vibration at pH 7 is due to the increased delocalization of ring electrons, because of greater electron density at the skeletal rings originating from conjugation with the carboxylate ion. However, the ID/IG ratio 0.94 for GO remains almost the same (0.91) in the hybrids of GO-­PVA. This is probably because of the ionized form of the carboxylic acid groups of GO at both in the pristine and hybrids at pH 7. However, at pH 4, the un-­ionized carboxylic acid group does not participate in delocalization of the ring electrons to the same extent, causing a shift of D bands to higher energy (1388 and 1469 cm−1). The doublet nature of the D band of GO at pH 4 may be attributed to the presence of two different types of sp3 skeletons present at GO; presumably one at the bulk, and the other at the side, the latter being severely passivated at pH 4, causing higher energy vibration. Using π-­stacking interaction Maity et al.22 prepared a rGO/PT-­g-­PMMA hybrid and studied the Raman spectra of the hybrid (Figure 4.19b) where rGO exhibits much narrower peaks of D and G bands from those of GO. It exhibits both D band at 1349 cm−1 and G band at 1588 cm−1 with the ID/IG ratio of 1.22. The higher ID/IG ratio of rGO from that of GO (0.97) arises from a decrease in the size of sp2 domains with an increase of its number for the

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restoration of the conjugated C=C bonds during the chemical reduction.31 In PFG, the ‘G’ band becomes broadened and shifts to 1585 cm−1 with a decrease of ID/IG ratio to 0.94 compared to rGO. This is for PT-­g-­PMMA with aromatic thiophene rings facilitating π–π interactions with the rGO assisting the G-­band vibration, thus easily lowering the G-­band position by 3 cm−1. This also decreases the ID/IG ratio from that of rGO possibly due to the increased planarity of rGO due to noncovalent grafting of the PT-­g-­PMMA chains on the rGO surface. So, Raman spectra suggest that the graphene sheets are highly functionalized with PT-­g-­PMMA by noncovalent π–π interactions. Layek et al.33 used pluronic F127 (PEO-­b-­PPO-­b-­PEO) surfactant to modify rGO (PF127-­rGO) and made a composite with poly(ethylene glycol), which was further used to plasticize gum arabic. They characterized PF127-­rGO using Raman spectroscopy to inspect the electronic structure of the carbon-­based materials (Figure 4.22). The Raman spectrum of GO exhibits two band positions at 1584 cm−1 (G-­band) and 1347 cm−1 (D-­band) and it is evident from the figure that both bands have shifted to lower energy vibration, 1580 cm−1 and 1343 cm−1, respectively, due to increased electronic conjugation for reduction of GO by hydrazine. The ID/IG ratio of GO (0.89) is found to be increased in rGO (1.07), possibly due to the formation of twisted rGO sheets due to its increased sp2 domains via the elimination of oxygen functional groups of GO. The surfactant functionalized

Figure 4.22  Raman  spectra of graphite, GO, rGO and PF127-­rGO. Reproduced from ref. 33 with permission from Elsevier, Copyright 2017.

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rGO (PF127-­rGO) shows a small shift in the G-­band to lower energy and the ID/IG ratio increases to 1.09, indicating that pluronic F127 has some influence, although small, possibly coming from nonbonding electrons of oxygen atoms present in the surfactant on the electronic environment of sp2 domains.

4.2.3  X-­ray Photoelectron Spectroscopy (XPS) XPS provides elemental composition, empirical formula, chemical and electronic states of the elements present within the material. The XPS spectrum of the GO/ATRP initiator exhibits a peak at 70.5 eV that corresponds to the binding energy of Br3d, confirming the initiator attachment on the GO surface. XPS spectrum is used to calculate the percentage of grafting of poly(dimethyl aminoethyl methacrylate) (PDMAEMA)from the area of N1s binding energy (399.8 eV) and C1s binding energy. The percentage of nitrogen calculated from XPS is ∼3.19 wt% in the grafted sample, yielding the PDMAEMA percentage, ∼35 wt%. The C1s spectrum of GO exhibits peaks at binding energies 284.6 eV (C–C), 285.7 eV (COH), 286.7 eV (C–O), 288 eV (C=O) and 289.1 eV (O–C=O).34,35 The C1s XPS spectrum of the sample shows all the C1s energy peaks but the percentage area of the peak at 289.1 eV (O–C=O) increased from 0.97% to 3.85%, indicating the grafting of PDMAEMA chains onto the GO sheets. The percentage of grafting of PDMAEMA, calculated from the increased area of C1s peak, is 32 wt%, close to that obtained from the above N1s area calculation. Qin et al.36 covalently functionalized GO using 1-­(3-­aminopropyl) pyrrole (APP) by esterification reaction between carboxyl and amino groups and subsequently in situ polymerized to produce GO/poly(1-­(3-­aminopropyl) pyrrole) copolymer. XPS was used to study the formation and the XPS wide scan spectra (Figure 4.23a and b) of GO shows no N1s peak; however, in the poly(GO– APP) sample a strong N1s peak is observed and the relative intensity of the C1s peak becomes enhanced for the grafted APP molecules. In the high-­ resolution XPS spectra (Figure 4.23c and d), the C1s spectrum for GO clearly shows five kinds of carbon atoms for C–C (284.6 eV), C–O (286.8 eV), >C=O (288.5 eV) and carboxylate O=C–O (290.2 eV) groups. However, the carboxyl groups (O=C–O) of GO sheets transforms into the amide groups (O=C–N) in poly(GO–APP). The C1s binding energy in the O=C–N groups (289.0 eV) is smaller than the O=C–O groups (290.1 eV). Also, the C=O peak shifted to a slightly lower binding energy (287.6 eV). This is because the electronegativity and inductive effect of N atoms are smaller than oxygen atoms causing the shift, giving clear evidence of the covalent functionalization of APP on the GO surface. X-­ray photoelectron spectroscopy (XPS) is also used to provide detailed information about the noncovalent functionalization of GO. Maity et al.22 noncovalently functionalized rGO with polythiophene-­graft-­poly(methyl methacrylate, PT-­g-­PMMA) (PFG), which has been subsequently used to make composites with poly(vinylidene fluoride). The XPS spectra (Figure 4.24a) portray the

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Figure 4.23  XPS  spectra of (a) GO and (b) a poly(GO–APP) copolymer (GO :  APP = 1 : 3.19), and high resolution C1s spectra of (c) GO and (d) a poly(GO–APP) copolymer (GO : APP = 1 : 3.19). Reproduced from ref. 36 with permission from the Royal Society of Chemistry.

characteristic peaks in the C1s, O1s and S2p regions of the rGO and PFG. The C1s core level spectra of rGO is analysed and is deconvoluted into four well-­ resolved Gaussian curves (Figure 4.24b) where the binding energy of four types of carbon appear at 285.9 eV for the C–C bond, at 287 eV for C–O of the epoxy, at 288.3 eV for C=O and at 289 eV for the carboxylate carbon (O=C–O).37 In PFG (Figure 4.24c) five types of carbon bond (C1s) with sp2 hybridized C=C (284.4 eV), sp3-­hybridized C–C (285.9 eV), C=O (288.3 eV), carboxyl groups O=C–O (289 eV) with the additional C–S bonds (285.7 eV), are present. Thus the XPS spectrum of the S2p region for PT-­g-­PMMA after surface functionalization is similar to different polythiophene derivatives.38 Deconvolution of the S2p (inset of Figure 4.24a) region indicates the existence of S2p3/2 (164.1 eV), and S2p1/2 (165.3 eV) peaks similar to sulfur of the polythiophene moiety38 supporting the noncovalent functionalization of PT-­g-­PMMA on the rGO surface. Gupta et al.39 used the γ-­radiolysis technique for the noncovalent functionalization of the rGO-­nanosheets (NSs) with poly(ethylene glycol) 200 (PEG200). The functionalization occurs via the intercalation of PEG200 molecules within rGO nanosheets through hydrogen bonding between the rGO and the PEG200 molecules. They used the XPS technique to investigate the chemical changes that occurred in the PEG-­rGO-­NS upon exposure to γ-­radiolysis. The broad peak in the C1s spectrum (Figure 4.25a(i)) is

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Figure 4.24  X-­  ray photoelectron spectra (XPS) of (a) GO and PFG (inset: S2p region of PFG). Deconvoluted spectral part of (b) C1s region of GO and (c) C1s region of PFG. Reproduced from ref. 22 with permission from Elsevier, Copyright 2016.

deconvoluted into three components, designated as A, B and C, at binding energies of 284.4, 285.9 and 288.6 eV, assigned to C–H/C–C, C–O–C/C–OH and C=O/HO–C=O carbon, respectively. The ratio of components A/(B + C) is 1.2 for PEG-­rGO-­NS, and is an important parameter to define the quantity of defects and to identify different functional groups present in the graphene structure. After γ-­radiolysis, the XPS spectrum of PEG-­rGONS shows significant change in the shape of the C1s core level XPS spectrum (Figure 4.25a(ii)). In the C1s spectrum, component A exhibits a band at 284.5 eV, and the B and C bands are observed at binding energies of 286.2 and 287.4 eV, respectively. The ratio of A/(B + C) intensity decreased from 1.2 to 0.92 which, together with the shifts of the A, B and C components from the pristine, indicate an interaction between rGO and PEG200 molecules. The sp2 fraction in PEG-­ rGO-­NS decreased from 70% to 62.2% on γ-­irradiation, pointing to the intercalation of –OH groups into the basal plane of the graphene sheets. This is a signature of PEG200-­functionalized graphene where functionalization is mediated through hydrogen bonding. The presence of phenolic groups and water molecules in graphene samples is clearly evident from O 1s core-­level

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Figure 4.25  XPS  spectra of (a(i) and b(i)) PEG-­rGO-­NS0 and (aii, bii) PEG-­rGONS3 in the (a(i) and a(ii)) C1s and (b(i) and b(ii)) O 1s regions. Reproduced from ref. 39 with permission from American Chemical Society, Copyright 2016.

spectra (Figure 4.25b(i) and (ii)) where, the three deconvoluted peaks A, B and C, observed at 530.1, 532.4 and 535.2 eV, respectively shift to 530.2, 532.5 and 534.5 eV, respectively, on irradiation (Figure 4.25b(ii)) providing evidence of H-­bonding interaction through the >C=O group and –OH group.

4.2.4  Wide Angle X-­ray Scattering (WAXS) WAXS is an important tool to diagnose the functionalization of graphene, which changes the exfoliation of graphene sheets. In WAXS spectra of graphite a diffraction peak at ∼2ϴ = 26.5°40 is attributed to the interlayer spacing of graphene sheets. However, for GO the diffraction peak shifts to a much lower angle at 2ϴ = 10.1° (Figure 4.26),41 suggesting that graphite has been successfully oxidized into GO, preventing the π–π stacking of graphene sheets, thus increasing the interlayer spacing. Feng et al.41 synthesized sulfonated graphene oxide (SGO) by direct esterification between graphene oxide (GO) and hydroxylated sulfonated poly(ether ether ketone) (SPEEK-­OH). The WAXS pattern of ester linked SGO shows a broad diffraction peak appearing at

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Figure 4.26  The  XRD curves of GO and SGO. Reproduced from ref. 41 with permission from Elsevier, Copyright 2018.

∼2ϴ = 24.58° indicating that the introduction of sulfonic acid groups increases the interfacial interaction between the graphene sheets and thus reduces the interlayer spacing between the GO layers42 suggesting the successful esterification of GO edge by the SPEEK-­OH. The WAXS patterns of poly(methyl methacrylate) functionalized GO (MG) by ATRP technique, poly(vinylidene fluoride) (PVDF) and the MG-­PVDF nanocomposites are presented in Figure 4.27.20 It is apparent that the MG has a small hump at 2ϴ = 8.2° indicating that the spacing between MG sheets increase from that of GO (cf. Figure 4.26). This is completely opposite to that of SGO, suggesting that there is no such interaction between MG sheets as in SGO. Also there is a big halo at 2ϴ = 23.8° in the MG spectra indicating MG is purely amorphous in nature. PVDF, being a semicrystalline polymer, exhibits α-­polymorphic crystalline diffraction peaks at 2ϴ = 17.8, 18.7, 20.3 and 26.8°.43,44 It is to be noted that the relative intensity of the 17.8, 18.8 and 26.8° peaks gradually decreases with increasing MG content in its composite with PVDF, suggesting a decrease in α-­polymorph and an increase in concentration of β-­polymorph having a characteristic peak at 2ϴ = 20.3° in the composite.43–46 It is interesting to note that the MG5 sample does not exhibit any characteristic diffraction peak of α-­PVDF and the β-­polymorph is produced fully. So, it may be concluded that MG induces β-­PVDF formation in the composite due to specific interaction with the >C=O group of MG with the >CF2 dipole of PVDF and the large surface area of graphene facilitating easier nucleation and growth of the β-­polymorph. Ding et al.14 synthesized GO using Hummer's method and reduced it thermally at 200 °C for 3 h to get TRGO from which a raft agent is grafted to get TRGO-­RAFT and finally polystyrene is grafted using AIBN as primary initiator. They used the XRD patterns (Figure 4.28) to understand whether graphene-­ based sheets exist as individual graphene sheets in the composites.

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Figure 4.27  WAXS  patterns of MG and different PVDF-­MG nanocomposites. Reproduced from ref. 13 with permission from Elsevier, Copyright 2010.

Figure 4.28  XRD  patterns of GO, TRGO, TRGO-­RAFT, PS and PRG. Reproduced from ref. 47 with permission from Elsevier, Copyright 2014.

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The 001 diffraction peak of GO is present at 2ϴ = 11°, with an interlayer spacing of 0.80 nm, suggesting the presence of oxygen-­containing functional groups.47 In TRGO, a broad peak at 24.4° is noticed, indicating the π-­stacked rGO sheets, and the smaller width of the peak of TRGO-­RAFT from that of TRGO signifies that the grain size of TRGO-­RAFT is larger than that of TRGO and the sheet layers become further exfoliated. The absence of a GO characteristic peak suggests that the layered TRGO is exfoliated in the composites.48 For the PS and PRG samples, two broad peaks are found at 2ϴ = 10 and 20° caused by amorphous PS, the first one is due to the intermolecular backbone–backbone correlation due to ordering of the molecular chains and the last one is due to the amorphous halo corresponding to the van der Waals distance.49 Maity et al.9 used the XRD technique to characterize noncovalently functionalized graphene with PT-­g-­PMMA by the π stacking process (Figure 4.29). The XRD pattern of graphite displays a single strong and sharp peak at 2ϴ = 26.45°, suggesting a well-­ordered lamellar structure having an interlayer d-­spacing value of 0.34 nm due to π-­stacking.50 After oxidation to GO a typical broad peak at 2ϴ = 10.26° for interlayer d001-­spacing (∼0.86 nm) is observed suggesting that GO sheets are well exfoliated. After reduction of GO, the peak of rGO exhibits a dramatic shift to higher 2ϴ = 23.92°, decreasing the d-­spacing value to 0.37 nm, which is slightly higher than that for bulk graphite denoting the reduction of GO. The lack of graphite peak indicates that the rGO sheets remain exfoliated and disorderly packed.51 PT-­g-­PMMA exhibits a broad peak at 2ϴ = 13.4° (dspacing = 0.66 nm) which is a feature of the inter-­stacking distance of PT chains. In PT-­g-­PMMA functionalized rGO (PFG) this peak shifts to 2ϴ = 14.4° (dspacing = 0.614 nm) and the small decrease in inter-­stacking distance in the hybrid may be due to its squeezing

Figure 4.29  XRD  spectra of graphite, GO, RGO, PT-­g-­PMMA and PFG respectively. Reproduced from ref. 9 with permission from Elsevier, Copyright 2016.

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by the cohesive force between the rGO nanosheets. Also the π-­stacking peak of rGO at 2ϴ = 23.92° has broadened suggesting that the PT-­g-­PMMA chains lie on the RGO sheet surface, thus disfavouring the inter-­stacking of rGO sheets. Gupta et al.39 used the XRD technique to functionalize rGO with polyethylene glycol (PEG 200) by hydrogen bonding between the oxygen atoms of PEG200 with the hydroxyl groups of rGO via γ-­irradiation. This PEG-­rGONS0 (without irradiation) exhibited a diffraction peak of the (002) plane at 2θ = 24.5° which shifted to 22.6° in the PEG-­rGO-­NS3 sample produced after 54 h irradiation. Thus, the interlayer d-­spacing of the (002) plane in PEG-­rGO-­NS0(0.36 nm) increased to 0.4 nm in PEG-­rGO-­NS3 (Figure 4.30). The increase in the interlayer spacing is attributed to the intercalation of PEG200 segments into the rGO matrix. The full width at half-­ maximum (fwhm) of the (002) plane of PEG-­rGO-­NS3 is narrower from that of PEG-­rGO-­NS0, suggesting an increase in the thickness of the graphene sheet along the c axis. The XRD data show that the total thickness of the graphene sheet is ∼2 nm in PEG-­rGO-­NS0, which increased to ∼3 nm in PEG-­rGO-­NS3. The diffraction peak from the (102) plane occurred at 2θ = 43 and 42.8° for PEG-­rGO-­NS0 and PEG-­rGO-­NS3, respectively. The fwhm of this peak, which is a measure of the lateral dimensions of the sheet, increased from 68 nm in PEG-­rGO-­NS0 to 82 nm in PEG-­rGO-­NS3 as a result of H-­bonding between the components at long time γ-­irradiation.

Figure 4.30  XRD  spectra of (a) PEG-­rGO-­NS0 and (b) PEG-­rGO-­NS3. Insets: typical interlayer spacings and sheet thicknesses. Reproduced from ref. 39 with permission form American Chemical Society, Copyright 2016.

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4.3  Thermal Properties The polymer functionalized graphene exhibits interesting thermal properties which include thermal stability, changing crystallization behaviour of the polymer-­like melting point, crystallinity, polymorph formation, etc. The thermal stability can be measured from thermogravimetric analysis (TGA) and the others can be measured from differential scanning calorimetry (DSC), which are discussed below.

4.3.1  Thermogravimetric Analysis (TGA) Under nitrogen atmosphere, TGA of GO loses its mass below 100 °C and here oxygen-­containing groups, e.g. –OH, epoxide, COOH, etc. degrade.51 However, the thermal stability of rGO sheets increases compared to that of GO for the elimination of a large fraction of oxygen-­containing groups during reduction.52 In PFG, more weight loss occurs for the degradation of the polymer grafted with graphene. The percentage of grafting of polymer from the graphene surfaces can be calculated using the difference between the weight losses of the two. By this method the amount of polystyrene attached by ATRP is found to be 82 wt% onto the graphene sheets.1 Yang et al.53 used the ATRP technique to graft PDMAEMA from GO and the composition of the nanocomposite was analysed by TGA. The exfoliated GO shows 13% weight loss in the range 100–800 °C for the loss of epoxy, –COOH and –OH groups on the GO sheets. The GO/PDMAEMA composite shows 32% weight loss upon grafting of the ATRP initiator and after grafting PDMAEMA the composites show 56% weight loss indicating that the weight percentage of PDMAEMA is ∼24%. Wang et al.54 grafted GO with hydrophilic poly(poly(ethylene glycol) methyl ether methacrylate) via the surface-­initiated ATRP (SI-­ATRP) method and have conducted TGA to investigate the thermal stability and composition of the synthesized nanoplates (Figure 4.31). A sharp weight loss between 150 and 300 °C for GO can be explained by the release of oxygen functionalities from the GO surface, and the sluggish weight loss between 300 and 800 °C is attributed to the pyrolysis of the other relatively stable oxygenous functional groups. Similarly, in GO-­NH2 and GO-­Br a weight loss of 61.7 and 63.7% is noticed, respectively. The different weight loss of GO (40.11%) and GO-­g-­P(PEGMA) (28.31–31.12%) between 100 and 450 °C indicate that the improved thermal stability of GO-­g-­P(PEGMA) at the above temperature is due to the presence of GO acting as heat barrier. On increasing the temperature to 800 °C, a relatively large loss in mass (%) is found due to the thermal decomposition of the grafted P(PEGMA) chains. The weight percentage of P(PEGMA) in GO-­g-­P(PEGMA) (0.5, 1 and 12 h), taking the weight loss of GO-­ Br at 800 °C as a reference, is about 5.1, 7.9 and 16.8%. The grafting amount of P(PEGMA) is calculated from an equation55 and is 5.6, 6.7 and 9.6 mmol g−1 for 0.5, 1 and 12 h, of polymerization. By using the esterification reaction technique Zhang et al.56 produced GO grafted hexadecane (GO-­HD) which is used to prepare GO-­HD/thermoplastic polyurethane (TPU) nanocomposites.

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Figure 4.31  TGA  curve of GO, GO–NH2, GO-­Br and GO-­g-­P(PEGMA) with different polymerization time. Reproduced from ref. 54 with permission from Elsevier, Copyright 2017.

Figure 4.32  Thermogravimetric  analysis data for GO and GO-­HD samples under nitrogen atmosphere with a temperature ramp rate of 10 °C min−1. Reproduced from ref. 56 with permission from Elsevier, Copyright 2016.

TGA was used to understand the thermal stabilities of the GO and GO-­HD samples (Figure 4.32). Apparently, GO shows a weight loss of ∼14% below 100 °C probably due to the evaporation of absorbed water molecules, followed by a significant weight loss (∼29 wt%) around 100–230 °C, resulting in the decomposition of labile oxygen-­containing functional groups. On the other hand, the GO-­HD sample exhibits only ∼1 wt% loss below 100 °C, indicating that the hydrophobicity of GO-­HD sample minimizes the amount of absorbed water greatly. The GO-­HD sample then suffers a dramatic weight loss (∼61%) at 100–400 °C

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due to the decomposition of the residual oxygen-­containing groups as well as long alkyl chains. All these results indicate that the long alkyl chains are successfully grafted onto GO surfaces through ester linkages. The derived content of grafted alkyl chains onto GO sheet surface is measured to be ∼28 wt%. Using the RAFT technique, Ding et al.14 grafted polystyrene (PS) chains on the GO surface by RAFT polymerization. In this procedure, a RAFT agent, 4-­Cyano-­4-­[(dodecylsulfanylthiocarbonyl) sulfanyl] pentanoic acid, is used to functionalize thermally reduced GO (TRGO) to obtain the precursor (TRGO-­ RAFT). TGA was used to understand thermal stability and the grafting degree of the RAFT agent and PS on the GO surface. Figure 4.33 shows that the weight loss of GO started below 100 °C due to the evaporation of absorbed water. Then, a relatively large weight loss was noticed at ∼200 °C, which was also true for TRGO and is ascribed to the decomposition of unstable oxygen-­ containing functional groups.57 The difference in the value of weight loss between TRGO and TRGO-­RAFT above 300 °C indicates the weight loss is for the decomposition of grafted RAFT agent. The weight loss curves for TRGO-­ RAFT, PS/graphene (PRG),and the neat PS, presented in Figure 4.33, clearly demonstrate that TRGO is effective in enhancing the thermal stability of PS. From the TGA curves of PRG and TRGO-­RAFT it is evident that the weight loss of PS decomposition is about 70 wt%, suggesting a high grafting quantity of PS chains on the GO surface. The grafting density can be measured from the formula according to TGA analysis47 as follow:

Functional groups per carbon: Amg =(MC − WF)/(MF − WC)

where MC is the relative molar mass of carbon (MC = 12 g mol−1), MF = molecular weight of functionalized groups of RAFT agent (MF = 403 g mol−1), and WC and WF are the weight fractions of the TRGO (excluding unstable groups of TRGO) and of the functionalized group, respectively. From the TGA curves the values of WC and WF are measured to be 80 and 20%, respectively.

Figure 4.33  TGA  curves of GO, TRGO, TRGO-­RAFT, PS and PRG. Reproduced from ref. 14 with permission from Elsevier, Copyright 2014.

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Then the grafting density of the RAFT agent on TRGO is calculated to be 0.84 functional groups per 100 carbon. In a similar way the grafting density of polymer chains are calculated from the following equation:

Polymer chains per carbon: = (MC − WP)/(MP − WC)

where MP is the number average molecular weight (Mn) of grafted polymer, and WC and WP are the weight fractions of the graphene backbone (excluding RAFT agent) and the grafted polymer, respectively. From the TGA curves one can easily obtain WC = 7% and WP = 93% in the PRG composites because the PRG shows a weight loss stage below 475 °C, and this is assigned to WP and the RAFT agent groups. Taking Mn of grafted PS from GPC = 86 922 g mol−1, the grafting density of PS chains on TRGO sheets is calculated to be 0.18 chains per 100 carbon. Noncovalent functionalization of rGO with polythiophene-­graft-­ poly(methyl methacrylate) (PT-­g-­PMMA) produced by π-­stacking interaction (PFG) was also characterized with TGA by Maity et al.9 to study the thermal stability and the amount of PT-­g-­PMMA functionalization (Figure 4.34a). GO has a low thermal stability and begins to degrade upon heating with weight loss of about 20% at ∼100 °C for the removal of intercalated water. Then another ∼33% weight loss occurs in the temperature range 100–250 °C for pyrolysis of the labile oxygen-­containing functional groups.58 In the case of rGO, the 20% mass loss below 100 °C is for the loss of adsorbed water into the π-­stacked graphene sheets and after that only 25% weight loss occurs at the higher temperature range till 800 °C. This development of thermal stability in rGO rather than that of GO is for the restoration of the graphitic structure in GO where oxygenous groups are mostly removed. In PT-­g-­ PMMA, the major weight loss starts at 250 °C showing weight loss of 91.8% at 440 °C. However, in PFG the major weight loss occurs at 250–440 °C showing 77.6% weight loss due to the disintegration of PT-­g-­PMMA on the rGO

Figure 4.34  (a)  TGA spectra of GO, RGO, PT-­g-­PMMA, PFG respectively in N2 atmosphere. (b) TGA curves of PRP composites and inset: variation of degradation temperature of PRP composites with wt% PFG. Reproduced from ref. 9 with permission from Elsevier, Copyright 2016.

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surface. Therefore, PFG contains 14.2 wt% of rGO, thus yielding a quantitative measure of noncovalent functionalization of PT-­g-­PMMA on the rGO surface by π-­stacking interaction. Utilizing the specific interaction between >C=O groups of PMMA in PFG and >CF2 dipole highly dispersed composites of PFG in PVDF are produced and are designated as PRP composites.9 TGA is used to understand thermal stability of PVDF in the PRP composites from the onset degradation temperature (Td) (Figure 4.34b). The Td of PVDF is 460 °C, but with increasing amount of PFG the Td at first increases to 470 °C up to 1% filler and then Td decreases. The improved thermal stability up to 1% filler is due to random dispersion of rGO sheet into the PVDF matrix for strong interfacial interaction between PVDF and the filler PFG that obstructs the heat flux required for degradation of the composite. Due to the good adhesion between PFG and PVDF, the thermal energies are transferred to the graphene, which causes an effective heat barrier. The high aspect ratio of graphene and better packing of PVDF chains in the composites hinder the onset of degradation. The sharp decrease in Td in the PFG3 and PFG5 samples is attributed to the self-­aggregation and change of orientation of the PFG sheets in these composites. Thus TGA offers a good idea about the functional characteristics and thermal stability of both covalently and noncovalently functionalized graphene.

4.3.2  Differential Scanning Calorimetry (DSC) DSC is also an important tool for understanding the effect of graphene on the melting, crystallization and polymorphic behaviour of graphene functionalized polymer and its composites. Layek et al. covalently functionalized rGO with PMMA by the ATRP technique and the modified graphene (MG) was used to make composites with PVDF through noncovalent interaction with the >C=O of MG and >CF2 dipole of PVDF. In Figure 4.35a, the DSC melting thermogram of melt cooled PVDF exhibits a single melting peak at 168 °C for the melting of α polymorph of PVDF. With addition of MG (0.5 wt%) another high temperature melting peak of PVDF appears at 175 °C which is attributed to the melting of β-­phase PVDF, and on increasing the MG concentration the melting peak of α-­phase PVDF decreases whereas that of β-­phase PVDF increases and in the 5 wt% MG composite only a β-­phase PVDF melting peak is noticed. To quantify the amount of α-­ and β-­phases of PVDF in the composite the DSC thermograms are deconvoluted and areas are measured (Figure 4.35b) where ∆H of α-­phase PVDF decreases with a concomitant rise in ∆H of β-­phase PVDF with increasing graphene content. In Figure 4.35c the DSC cooling thermograms of the composite samples are presented where a gradual increase in the crystallization temperature with increase in graphene content increases. So, graphene sheets act as a nucleating agent by providing its very large surface area for adsorption of PVDF chains and thereby causing easier nucleation. The increase in β-­phase content with MG concentration may be attributed to the specific interaction between the >C=O group of PMMA in MG with the >CF2

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Figure 4.35  (a)  DSC melting endotherms (scan rate 40 °C min−1) of different melt-­ cooled PVDF-­MG nanocomposites. (b) Plot of enthalpy of α-­ and β-­phases with MG content (w/w %) (obtained from thermograms in (a)). (c) DSC cooling endotherms (scan rate 5 °C min−1) of PVDF and different PVDF-­MG nanocomposites. Reproduced from ref. 13 with permission from Elsevier, Copyright 2010.

group of PVDF facilitating all-­trans conformation of the β-­phase assisted by the large surface area of graphene. Maity et al.59 covalently functionalized GO with imidazolium ionic liquid (IL) and the resultant functionalized graphene oxide (GO-­IL) was used to produce composites with PVDF (PGL) by the solution casting method. The melting and crystallization thermograms studied by DSC are presented for pure PVDF and its PGL composites in Figure 4.36a–c. Pure PVDF shows a melting peak at 168.7 °C for its α crystals, and this area decreases in the PGL composite showing the appearance of a new peak at ∼171.7 °C due to melting of the ß-­phase.60,61 The shifting of the 168.7 °C peak of the α-­phase of

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Figure 4.36  (a)  DSC melting endotherms of PVDF and the PGL composites at indi-

cated compositions at 10 °C min−1. (b) Plot of ΔHα and ΔHβ phases with GO-­IL content (w/w %) (obtained from heating thermograms). (c) Crystallization isotherms of PVDF and different PGL composites films for cooling from the melt at the rate of 5 °C min−1. Reproduced from ref. 59 with permission from Elsevier, Copyright 2015.

pure PVDF to a lower temperature in the composite is due to its supramolecular interaction with GO-­IL.62 So, the ionic liquid in GO-­IL acts as a good interfacial medium between PVDF and the graphene sheet, thus playing an important role as a compatibilizer facilitating homogeneous dispersion of graphene in the PVDF matrix facilitating β-­phase PVDF formation.13,63 To get an approximate idea about the quantity of α-­ and β-­phases in PGL composites the melting peaks of α-­and ß-­phases are deconvoluted and the enthalpy of fusion of both the phases are computed from the respective areas. Figure 4.36b indicates a sharp decrease in the α-­phase enthalpy of fusion and a sigmoidal rise in the case of the ß-­phase. The figure shows that in the PGL3 composite the α-­phase enthalpy of fusion decreases to only 10 J g−1 and this residual α-­phase arises for the nucleation far from the GO-­IL surface, i.e. at the bulk of the composite. The cooling thermograms (Figure 4.36c) of PVDF shows the crystallization temperature (Tc) at 134.7 °C and it gradually shifts

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to higher temperatures in the composites. This result indicates that GO-­IL, for its large surface area, acts as a nucleating agent, accelerating the crystallization of PVDF and also inducing the formation of β-­crystals from the melt. This synergetic effect is ascribed to the supramolecular interaction between the PVDF chains and the GO-­IL.64,65 Song et al.66 functionalized GO with pyrene-­functionalized poly(methyl methacrylate)-­block-­polydimethylsiloxane (Py-­PMMA-­b-­PDMS) by noncovalent π–π interaction between pyrene and GO sheets and measured the glass transition temperature (Tg) of PMMA from the DSC curves. Tg of PMMA decreases from 116.6 to 103.6 °C on addition of 0.5 wt% Py-­PMMA-­b-­PDMS. This result suggests that Py-­PMMA-­b-­PDMS has a plasticizing effect owing to the flexible Si–O–Si segments and good dispersion for the same PMMA segments in Py-­PMMA-­b-­PDMS with the PMMA matrix. The glass transition temperature (Tg) of all the composites gradually decreases with increasing GO concentration, demonstrating that high loading of GO affords more active sites to interact with Py-­PMMA-­b-­PDMS by π–π interaction, resulting in more uncoiled Py-­PMMA-­b-­PDMS chains. Thus, Py-­PMMA-­b-­PDMS acts as a good compatibilizer between GO and PMMA matrix in the composite. Maity et al.22 utilized π–π interaction for noncovalent functionalization between polythiophene-­graft-­poly(methyl methacrylate) (PT-­g-­PMMA) with rGO (PFG) which is used as a filler to make composites with PVDF (PRPx, where x indicates PFG content (wt%)). rGO sheets become distributed homogeneously in the PVDF matrix due to supramolecular (dipolar) interaction between the >C=O group of PMMA and >CF2 groups of PVDF chains. In Figure 4.37a the DSC melting thermograms of PVDF and PRP composite films indicate that PVDF exhibits one melting peak for the α-­phase and the PRP composites (PRP0.1–PRP3) exhibit two melting peaks, the lower melting

Figure 4.37  (a)  DSC melting endotherms of PVDF and the PRP composites at indicated wt% of PFG at the heating rate of 40 °C min−1. (b) Variation of melting temperatures of β form Tβ, α form Tα and crystallization temperature Tc of PVDF and PRP composites. Reproduced from ref. 22 with permission from Elsevier, Copyright 2016.

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peaks of PRP0.1– PRP3 are for the α-­phase whereas the higher melting peak is for the β-­phase.61,67 The melting points of both the phases with composition are presented in Figure 4.37b and it is evident from the figure that the melting points of the α-­phase show initially a fast decrease and then a slow decrease. This is due to (i) interaction between grafted PMMA in PFG and PVDF chains at the melt68 and (ii) formation of thinner α-­phase crystal for the growth of β-­phase crystal. Also the melting peak of the β-­phase increases gradually with increase in PFG concentration and maybe for its increased crystallinity and increased crystal thickness due to creation of a major fraction of the β-­phase. However, total PVDF crystallinity decreases due to the confinement of some fraction of PVDF chains at the interphases of the functionalized graphene.69,70 The cooling thermograms of PVDF and PRP composites indicates that there is a gradual increase in crystallization temperature with increase of PRP concentration indicating acceleration of PVDF crystallization, which is due to the strong heterogeneous nucleation of PVDF on the large graphene surface of the composites.

4.4  Conclusion It is very much evident from the above discussion that the morphology, structure, and spectral and thermal properties of both polymer and graphene become altered in the presence of GO/rGO in PFG. The morphology correctly infers the functionalization of graphene, spectral properties characterize the interaction between the components, WAXS and XPS characterize the structural changes due to interaction of polymer due to the presence of filler, and the thermal properties give an idea about the thermal stability, polymorphic nature of the polymer in PFG, etc. In a word, graphene alters all the above physical properties of the polymer due to synergistic interaction between the filler and matrix polymer.

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Chapter 5

Optical Properties of Polymer Functionalized Graphene: Application as Optical Sensor 5.1  Introduction Graphite is a black material, hence it absorbs all visible wavelengths of light; however, a thin layer of graphene is optically transparent. The continuous sp2-­hybridized structure of graphene makes it optically inactive but GO possesses ∼50% of sp3-­hybridized carbon containing different oxygenated functional groups producing a localized sp2 structure. Thus both π–π* and n–π* absorption bands are expected in the UV–vis spectra. Also the localized sp2 structure creates a band gap, which may yield fluorescence due to nonradiative decay of excitons in the localized states of graphene. Hence fluorescence is observed in GO solution at acidic pH and is also observed in polymer functionalized graphene (PFG). This fluorescence property is very useful for fabricating chemical and biochemical sensors and in this chapter we shall delineate the optical properties (UV–vis and fluorescence) in details including the p-­and n-­t ype doping of graphene in PFG with some examples of optical sensors.

5.2  UV–Vis Spectra Graphene has a continuous sp2-­hybridized structure while GO possesses a large amount (∼50%) of sp3-­hybridized carbon containing different oxygenated functional groups. The UV–vis absorption spectra of pure GO is   Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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presented in Figure 5.1a and it shows absorption peaks at 240 and 304 nm attributed to π–π* and n–π* transitions, respectively.1 The n–π* transition arises from the transition of nonbonding electrons of oxygenated groups of GO to its π* orbital at its lowest excited state. In Figure 5.1a the change in UV–vis spectra on adding methyl cellulose (MC) is presented.2 The π–π* transition peak at 240 nm shows a blue shift to 235 nm in the GMC hybrids, possibly due to a decrease in planarity of GO arising from H-­bonding with MC forming a lesser conjugated system. The n–π* absorption peak vanishes for the H-­bond formation through the nonbonding electrons of oxygen atoms of GO with the hydroxyl groups of MC. In addition, the absorbance spectra exhibit a long tail which indicates that the π–π* transition in GO can occur anywhere between 220 and 600 nm.3 Figure 5.1b shows the UV–vis spectra of GO and GO-­ poly(vinyl alcohol) (PVA) solutions of different concentrations at pH 4 where pure GO exhibits two absorption peaks at 244 and 302 nm (inset) for the π–π* and n–π* transitions, respectively.4 Thus there is a small effect on the absorption peak positions of GO

Figure 5.1  (a)  UV–vis spectra of GO and GMC hybrids solutions at pH 7. Reproduced from ref. 2. (b) UV–vis spectra of GO and GO-­PVA hybrids of different compositions at pH 4. Reproduced from ref. 4 with permission from American Chemical Society, Copyright 2012.

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with pH; the former peak becomes red shifted by 4 nm while the latter is blue shifted by 2 nm. This is because at pH 7 the carboxylic acid groups remain ionized while at pH 4 it becomes unionized, affecting the electronic density on the graphene ring and carboxyl groups. The UV–vis spectrum of GO at pH 4 changes on addition of PVA, causing blue shift in both the absorption peaks. The 244 nm peak becomes shifted to 210, 208, and 204 nm for the GO-­PVA0.5, GOPVA1, and GO-­PVA2 samples, respectively, while the 302 nm peak has shifted to 258 nm in every case. The former blue shifts increase with increase in PVA concentration and is attributed to the decrease in planarity of GO sheets due to hydrogen bonding with the high-­molecular-­weight PVA chain,5,6 forming a lesser-­conjugated cranked graphene sheet, thus increasing the band gap. Ji et al.7 used pyrene-­terminated liquid crystalline polymer (LCP, BP6-­6400 g mol−1) to graft GO, exploiting the π–π stacking interaction, which is supported from UV–vis spectra in DMF (Figure 5.2). The GO spectrum in water exhibits two peaks at ∼230 and ∼300 nm (Figure 5.2A), corresponding to the π–π* transitions of the aromatic C=C band and n–π* transitions of the C=O band of GO, respectively. However, the spectra of GO in DMF shows only a shoulder at ∼300 nm indicating that GO can be dispersed in DMF. Figure 5.2B shows the sharp main absorption peak of LCPs at 297 nm and this is true for LCP functionalized graphene sheets, indicating the existence of LCP on the graphene surface. However, these spectra of hybrids become broader around 297 nm from pristine LCP due to the presence of π-­stacking interactions between the graphene sheets and pyrene functional groups attached to the LCPs. Thus, shifting the π–π* and n–π* absorption peaks of GO or broadening or abolition of these peaks in UV–vis spectra clearly indicates noncovalent interaction between GO and the polymer.

Figure 5.2  UV–vis  absorption spectra of (A) GO aqueous solution and (B) dispersion of LCP and LCP functionalized graphene sheets in DMF (insert: GO in DMF). Reproduced from ref. 7 with permission from Elsevier, Copyright 2015.

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5.3  Photoluminescence (PL) Spectra Due to the presence of a continuous sp2-­hybridized zero band gap structure graphene does not exhibit any fluorescence properties. However, due to the presence of a large amount (∼50%) of sp3-­hybridized carbon, GO possesses a finite electronic band gap.8–10 This finite band gap of GO causes PL properties in the visible and near-­infrared region.11 The presence of a localized molecular sp2 cluster within a sp3 matrix confines the π electrons and the size of the sp2 cluster portrays the local energy gap.12 The radiative recombination of localized electron–hole (e–h) pairs in sp2 clusters generates the PL properties13,14 and by tuning the size, shape and fraction of sp2 clusters, the PL intensity may be controlled.12,15 Gokus et al. noticed visible luminescence from oxygen plasma treated graphene and characterized it to the electron confinement in sp2 islands of ∼1 nm size.16 The poor emissive property of GO originates from the hydroxyl, and epoxy groups which induce nonradiative recombination by the transfer of their nonbonding electrons to the holes.11,15 However, the partial reduction of GO causes enhancement of PL intensity due to the formation of more localized sp2 clusters and structural defects.5,15 The reduction of a thin film of GO with hydrazine hydrate vapour for a short time period (0.4% w/v) at PH 4 gelation occurs due to the aggregation of GO sheets.19 However, at much lower GO concentration (0.005% w/v), gelation does not occur, but here the protonation of carboxyl and epoxy groups decreases the non-­radiative recombination of the localized e–h pair significantly, causing a large increase in emission intensity. Here, some interlayer quenching within GO sheets can occur due to aggregation; however, it is much lower from the quenching with the smaller size solvent (water) molecules, contributing to the increased emission intensity. At higher pH more ionization of carboxylic acid groups cause increase of non-­radiative recombination of

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the e–h pair thereby decreasing the PL intensity to one-­third to that of the neutral medium.

5.3.1  Fluorescence in Polymer Functionalized Graphene It would now be exciting to discuss if it is possible to passivate the GO surface more efficiently by decreasing the electron donating power of hydroxyl, epoxy and carboxylic acid groups of GO to the holes of graphene rings produced at the excited state by making H-­bonds with suitable polymers. Thus, the non-­ radiative (e–h) recombination centres can be decreased, which can improve the PL property dramatically. This property can be used for sensing of different molecules, ions and biomolecules. Further, the functionalized polymer can dope graphene, altering the electron density differently for change of pH, consequently a change in fluorescence properties occurs. We shall discuss them in the following sections taking individual PFGs produced by both noncovalent and covalent functionalization.

5.3.1.1 Methyl Cellulose Functionalized GO in Enhancing Fluorescence Nandi and co-­workers2 chose the formation of a hybrid of methyl cellulose (MC) with GO by mixing it in GO solution. The hybrids are designated as GMCx, where x denotes the MC concentration (w/v) in the solution.2 In Figure 5.3b the fluorescence spectra of GO at pH 4 and that of the GO–MC hybrid at different composition are shown and it is evident that the PL intensity augments significantly with increasing MC concentration. At pH 7 and 9.2, although small, an increase in PL intensity is also noticed with MC concentration (Figure 5.3c). The increase in PL intensity with MC concentration is compared for different pH in Figure 5.3c, which shows that the PL-­intensity at pH 4 rises by 7.8, 4.0 and 2.3 times over pH 7 for the addition of 0.85, 1.7 and 3.4% (w/v) MC solutions, respectively. The reason for the increase in PL intensity of GO at pH 4 compared with that of pH 7 is the protonation of carboxylate ion of GO19 and this is also true for the GMC0.85, GMC1.7 and GMC3.4 hybrid solutions. This is because the hydroxyl groups of MC make H-­bonds with the –COOH, –OH and epoxy groups of GO, decreasing the nonradiative recombination of the electron–hole pair causing enhancement of PL intensity. The higher molecular weight of MC is a possible reason for its preferential H-­bonding with GO than with water molecules, thus the H-­bonded MC avoids the quenching of excitons with the water molecules. This is also supported by the increase in PL intensity with increase in MC concentration, where the increased H-­bonds between GO and MC augment the passivity of the GO surface, causing an enhancement of emission efficiency. The unaltered position of emission peak at 437 nm with change in pH or with addition of MC suggests that the average cluster size of sp2 domains remains almost unchanged during the processes.

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Figure 5.3  (a)  PL spectra of GO at different pH. (b) PL spectra of GMC at differ-

ent concentrations at pH 4. (c) PL-­intensity versus MC concentration at different pH values. Reproduced from ref. 2 with permission from the Royal Society of Chemistry.

The lifetime values of GO at pH 4 and GMC3.4 at different pH values are measured and the average lifetime values along with the component decay times are presented in Table 5.1. It is evident from the table that the average lifetime of GO at pH 4 increases in the GMC3.4 hybrid, indicating the formation of a more stable excited state. This is possibly due to the H-­bonding between the hydroxyl group of MC with different functional groups of GO. On increasing pH, the lifetime values of GMC3.4 also increases indicating

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Table 5.1  Lifetime  values of GO at pH = 4 and GMC3.4 at pH = 4, 7 and 9.2 (ref. 2).

Systems

Average Relative Relative Relative Relative life amplitude τ2 amplitude τ3 amplitude τ4 amplitude time τ1 (ns) (ns) (a1) (ns) (a2) (ns) (a3) (ns) (a4)

GO(pH 4) 0.05 GMC 0.94 (pH 4) GMC 1.09 (pH 7) GMC 1.03 (pH 9.2)

15.02 16.84

1.15 4.94 3.91 46.11

3.78 70.13 9.17 19.24

7.17 9.91 0.05 17.81

3.42 3.73

28.46

4.14 31.65

12.2 21.78

0.11 18.10

4.30

23.08

4.02 34.29

13.2 27.10

0.10 15.52

5.21

that the excitons become more stabilized than those at pH 4 although the PL intensity has considerably decreased (Figure 5.3c). This suggests that at these high pH values (pH 7 and 9.2), a less fluorescent complex is produced at the ground state20 due to the formation of carboxylate and phenolate ions on the GO surface facilitating the nonradiative recombination of the electrons of ions to the nearby holes. This causes a decrease in radiative recombination of excitons with the holes and exciton decay through an alternate path. 5.3.1.1.1  Sensing of Picric Acid by GO–MC.  Being interested in its possible application as fluorescent sensor, the changes in PL intensity of GMC0.85 for adding different aromatic compounds at pH 4 is tested. On adding benzoic acid (BA), phenol (PH) and aniline (AN) (170 µM each) a small decrease in PL intensity occurs, but on addition of the same concentration of nitroaromatics, such as nitrophenol, 2,4-­dinitrophenol (DNP) and picric acid, significant quenching of PL intensity occurs (Figure 5.4a). On adding 170 µM picric acid solution a significant quenching of PL intensity occurs with a gradual red shift of the emission peak on increasing picric acid concentration. The nitroaromatics are electron deficient centres (hole-­like) facilitating exciton transfer from the electron rich sp2 cluster of the GO–MC hybrid, hence reducing the radiative recombination of the electron–hole pair, causing significant quenching of PL intensity. The quenching of fluorescence intensity is instantaneous, attaining the minimum value in C=O group of the block copolymer breaks, making GO more planar, thus decreasing the band gap causing the red shift to 459 nm. The intensity of 410 nm peak shows a large decrease (almost flat) when the temperature is increased to 35 °C. This peak originates from the n-­t ype doping of GO and it becomes disfavoured due to the onset of macro phase separation at 35 °C. With more increase in the temperature by 10 °C, decoiling of the polymer chain begins and doping from the –NMe2 groups starts again, showing an increase in the intensity of the 410 nm peak. However, the intensity of the 455 nm peak decreases with increase in the temperature,

Figure 5.14  Solid  state PL spectra of GPCLD at pH 9.2. Reproduced from ref. 41 with permission from American Chemical Society, Copyright 2017.

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as the aggregation of grafted chains facilitates the graphene sheets to be nearer causing nonradiative decay of excitons, hence decreasing the fluorescence intensity.

5.3.5  F  luorescent Amylose-­functionalized Graphene: Chiral Detection Wei et al.43 developed a reusable natural cheap polysaccharide (amylose) functionalized graphene for highly sensitive and visual fluorescent chiral sensing of the biological molecule, tryptophan (Trp). They first synthesized the fluorescent dye anthracene-­labelled amylose (AA) and then noncovalently functionalized it with rGO to produce AA–rGO hybrids exploiting the interactions between rGO and amylose chains.44,45 On incubation of rGO with AA it undergoes stacking interactions and also produces H-­bonding with oxygen residues via hydroxyl groups. The resulting AA–rGO is water soluble and nonfluorescent due to strong energy transfer between the fluorescent anthracene moiety and graphene. Upon exposure of AA–rGO to target molecules, AA molecules are released from rGO and make an AA-­target complex causing the recovery of anthracene fluorescence. In Figure 5.15 the fluorescence spectrum of AA in PBS in absence of rGO exhibits strong fluorescence (curve a); however, for the AA-­rGO hybrid, about 98.5% fluorescence becomes quenched (curve c). This result suggests adsorption of AA on rGO causing high fluorescence quenching efficiency. Intriguingly, upon addition of L-­/D-­Trp, the AA–rGO hybrids displayed chiral-­selective fluorescence enhancement (curves d and e). The fluorescence intensity of free AA is not quenched by the addition of L-­Trp (curve b). In the inset of Figure 5.15 the fluorescence intensity ratio (F/F0) of AA and AA–rGO is compared for the addition of L-­Trp, where F0 and F are the fluorescence intensity at 500 nm in the absence or presence of L-­Trp, respectively. In the case of pure AA, fluorescence has hardly changed for adding a high concentration of L-­Trp. However, in the AA–rGO hybrids, a 70-­fold increase in fluorescence intensity is noticed. DMSO, having a low dielectric constant (47.2), is known to induce a conformational transition of amylose to helical structure by increasing the hydrogen bonds between adjacent sugar and the fluorescence intensity of AA–rGO hybrids become gradually enhanced with increasing DMSO. However, the fluorescence intensity of AA alone does not change with increasing DMSO suggesting that when DMSO is added to AA–rGO solutions, AA forms inclusion complex like helical structures causing release of AA from rGO, thus causing recovery of fluorescence. Of particular importance for a chiral sensor is its chiral selectivity against the enantiomers and for this purpose the ratio of the fluorescence enhancement of the G-­based chiral sensor against D-­/L-­Trp (FL/FD) is compared. It is evident from Figure 5.16a that L-­Trp binding causes significant fluorescence enhancement, whereas a much lower fluorescence increase is noticed for D-­Trp binding at the same concentration (75 mM).

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Figure 5.15  Fluorescence  emission spectra of AA (5 mg mL−1) at different condi-

tions: (a) AA in PBS; (b) AA + 500 mM L-­Trp; (c) AA-­rGO; (d) AA rGO + 500 mM L-­Trp; (e) AA-­rGO + 500 mM D-­Trp. All the measurements were carried out in 25 mM PBS (pH 6). The AA–rGO concentration was kept at 10 mg mL−1. Inset: fluorescence intensity ratio of AA–rGO and the control (AA) with 500 mM L-­Trp. Excitation: 400 nm; emission: 500 nm. Reproduced from ref. 43 with permission from the Royal Society of Chemistry.

The chiral selectivity (F-­L/F-­D) of this rGO–AA sensor toward Trp enantiomers is 3.7, indicating that the chiral selectivity of a rGO–AA sensor is better than recently reported chiral sensors.46,47 Further, the chiral sensing and the concomitant fluorescence difference can also be recognized under a UV lamp. For the same concentration of D-­and L-­Trp enantiomers the colour of the rGO-­AA solutions are blue and green (inset of Figure 5.16a), respectively because a 14 nm red shift (489 to 503 nm) of the fluorescence peak occurs for the blue to green colour change, and such fluorescence colour difference can be readily detected by the naked eye. Time-­dependent experiments show that the fluorescence enhancement of the rGO–AA sensor against L-­Trp occurs rapidly in the first 2 min, and then the increase is sluggish over the next 6 min whereby the sensing is complete within 10 min, suggesting it as a fast process. The selectivity of rGO–AA sensing for Trp toward other essential amino acids is shown in Figure 5.16b. The other two aromatic amino acids, Tyr and Phe, assemble to AA much weaker than Trp, hence do not hinder in the selective sensing of Trp. Very high concentrations (10 mM) of nonaromatic amino acids, including arginine (Arg), cysteine (Cys), serine (Ser), isoleucine (Ile), threonine (Thr), and histidine (His), can only induce little fluorescence recovery of AA, indicating a strong anti-­interference ability of the rGO–AA sensor towards these nonaromatic amino acids. Trp selectivity arises from

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Figure 5.16  (a)  Fluorescence emission spectra in of AA-­rGO in the absence and

presence of 75 mM D-­or L-­Trp. Inset: a photograph for AA-­rGO with 75 mM D-­, or L-­Trp excited by a hand-­held UV lamp. (b) Selectivity of the G-­based platform for L-­, D-­Trp toward other essential amino acids. L-­Trp, D-­Trp, L-­Tyr, D-­Tyr, and L-­Phe were tested at 100 mM, and other nonaromatic amino acids were tested at 10 mM. Reproduced from ref. 43 with permission from the Royal Society of Chemistry.

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48,49

its specific interaction with amylose and it remains at the active sites of glucoamylase, inducing a conformational change in amylose and loses its enzymatic activity towards its substrate amylose when they are mutated by other amino acids. The results support the specific selectivity of the G-­based platform for Trp towards other amino acids. This specific selectivity for tryptophan (Trp) enantiomers amongst other essential amino acids permits potential chiroselective analysis of Trp in biological fluids with rGO–AA and, in principle, the present design can be used for sensing other target molecules that can produce an inclusion complex with amylose.

5.3.6  β  -­Cyclodextrin Functionalized Graphene: Fluorescent Detection of Cholesterol Mondal and Jana50 intriguingly developed a fluorescence based cholesterol detection method using competitive host–guest interaction between graphene grafted β-­cyclodextrin (β-­CD) with rhodamine 6G (R6G) and cholesterol. They synthesized a water soluble β-­CD–G by in situ reduction of GO in the presence of β-­CD. When GO is reduced in the presence of β-­CD, it becomes highly water soluble, even making it free from unbound β-­CD via extensive dialysis. β-­CD becomes covalently attached to the GO surface via reaction of the OH group present in β-­CD with the epoxide groups of the GO surface,51 producing β-­CD–G. The β-­CD–G displays the property of both components; β-­CD acts as a host for different guest molecules and graphene acts as a strong fluorescence quencher for the guest fluorophore rhodamine 6G (R6G). The absorption spectra show characteristic bands at 520 nm for R6G and at 260 nm for graphene. The absorption peak of R6G shows red shift of 25 nm from that of free R6G, due to interaction with graphene and its fluorescence is drastically quenched due to charge and energy transfer.52 The lifetime of R6G excitons increases once it enters into the β-­CD host but decreases below the nanosecond range in β-­CD–G, which further proves the incorporation of R6G into the β-­CD–G, thus graphene attached to β-­CD acts as a fluorescence quencher. The R6G present inside the β-­CD–G can be selectively replaced by cholesterol because of its high binding affinity to the β-­CD cavity due to its hydrophobic nature. This replacement of R6G by cholesterol releases R6G in the solution resulting fluorescence ‘turning on’ (Figure 5.17a) and the increased fluorescence is directly proportional to the amount of cholesterol added. The observed fluorescence response can be used for naked eye based semiquantitative detection and spectrometer based quantitative detection of cholesterol and the detection sensitivity can reach up to the nanomolar concentration range (Figure 5.17b) which is comparable to most of the existing cholesterol detection methods.53 The mechanism of fluorescence sensing of cholesterol is clearly illustrated in Figure 5.17c. The method is good, showing no interference from common molecules present in human blood serum and also in different salts, carbohydrates and amino acids. Thus this method is successful for determining cholesterol in blood.

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Figure 5.17  Cholesterol  induced ‘turn on’ fluorescence: (a) different amounts of

cholesterol are added into a solution of β-­CD–G–R6G. The inset shows a cholesterol concentration dependent fluorescence intensity plot with a fitting curve. (b) Digital image of a solution of β-­CD–G–R6G before and after cholesterol addition. (c) Schematic presentation of fluorescence based cholesterol detection using β-­CD–G via competitive host–guest interaction. Fluorescence of R6G inside β-­CD is quenched by graphene but it ‘turns on’ as cholesterol replaces R6G. Reproduced from ref. 50 with permission from the Royal Society of Chemistry.

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5.3.7  F  luorescent Block Copolymer-­functionalized Graphene Oxide: Efficient Temperature Sensing Yang et al.54 reported a competent temperature sensing platform from GO grafted with a fluorescent, thermally responsive block copolymer. They have synthesized P7AC-­b-­PNIPAM-­b-­PSN3 triblock copolymer consisting of poly(7-­ (4-­(acryloyloxy)butoxy)coumarin) (P7AC) as the fluorescent component, poly(N-­isopropylacrylamide) (PNIPAM) as the thermally responsive polymer, and a short poly(azidostyrene) (PSN3) block by a RAFT polymerization technique (Figure 5.18). The azide functional groups in PSN3 facilitate covalently linking P7AC-­b-­ PNIPAM to the GO surface producing fluorescent, thermally responsive GO composites (FGO). The thermal response of FGO in water is determined by monitoring the change in fluorescent intensity with temperature. The FGO composites dispersed in water exhibited a distinct blue colour (λ ≈ 425 nm) under irradiation with a UV lamp (365 nm). They studied the temperature dependent fluorescence of both FGO1 and FGO2 attached with two different molecular weight (Mn) values of PNIPAM. When the temperature of the solution is above 32 °C (LCST of PNIPAM), the fluorescence intensity becomes reduced dramatically for both FGO1 and FGO2 (Figure 5.19a and b). In a control experiment, the fluorescence intensity of ungrafted P7AC-­b-­PNIPAM-­b-­ PSN3 copolymers in water with increasing temperature does not show any

Figure 5.18  Schematic  illustration of the FGO and its application as a thermal sensor Reproduced from ref. 54 with permission from the Royal Society of Chemistry.

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Figure 5.19  Photoluminescence  (PL) spectra of (a) FGO1 and (b) FGO2. (c) Changes

in PL intensity at 425 nm as a function of temperature. Reproduced from ref. 54 with permission from the Royal Society of Chemistry.

change. So, the dramatic change in fluorescence intensity of FGO1 and FGO2 arises from the FRET efficiency between the P7AC block and GO, as the conformation of PNIPAM changes with increasing temperature. Below the LCST, PNIPAM is hydrophilic with expanded chain conformation, thus the distance between the P7AC and the GO surface is higher, thereby suppressing FRET, but above the LCST, the distance between the P7AC and the GO surface is lower. The LCST of 32 °C for the coil-­to-­globule transition of PNIPAM is consistent with the fluorescence intensity changes as a function of temperature (Figure 5.19c). FRET from the blue-­emitting coumarins in P7AC block and GO becomes controlled by its interspacing, which is governed through the manipulation of conformational features of PNIPAM chains affected by temperature changes. Interestingly, the PL quenching efficiency of FGOs depends on the Mn value of the PNIPAM spacer of P7AC-­b-­PNIPAM-­b-­PSN3 chains. After heating FGO1 in water above 32 °C, the maximum PL intensity decreased by 39%, while the intensity of FGO2 decreased by 44%. The molecular weight of PNIPAMs in FGO1 and FGO2 are 15 K Da and 33 K Da, respectively. So, the higher PL intensity decrease in FGO2 than FGO1 is because of the larger globule state of PNIPAM that provides appropriate distance between the GO and the P7AC block commensurate for efficient quenching.

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5.4  Scope In conclusion, we may say that the optical properties of PFG are very much interesting as pristine graphene is a zero band gap material. Polymer functionalization via the GO route induces a new band gap that facilitates obtaining both absorbance and fluorescence spectra. Depending on the nature of polymer grafted on the GO surface, a variety of fluorescence properties can be generated which depends on temperature, pH, etc. The pH dependent fluorescence not only causes sensing of different molecules but also gives an idea about p-­and n-­t ype doping. In a wider sense we can get an idea of localized and delocalized doping when block copolymers are grafted. A suitable choice of block-­copolymer having a component exhibiting LCST can also find use in temperature sensing. The fluorescence properties of PFG are very useful in sensing nitroaromatics, metal ion, chiral molecules, chlorosterol, etc. So the scope of optical properties of PFGs are very important for fabricating different types of sensors and other optoelectronic applications.

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Chapter 6

Mechanical Properties of Polymer Functionalized Graphene 6.1  Introduction Among the many unique characteristics of nanomaterials, special attention may be drawn to their large surface-­to-­volume ratios and extraordinary mechanical properties. These properties offer avenues for stimulating areas of research and also for technological innovations. So, a main use of nanomaterials is in reinforcing polymer matrices exploiting their ultra-­high stiffness and hardness. Recent research has shown that small additions (∼1 wt%) of certain nanomaterials can augment the mechanical properties abruptly, sometimes by as much as 100%.1–6 The precise mechanism for this dramatic enhancement is not clearly understood and it is thought that molecular level interactions between nanomaterials and polymer matrices play a crucial role. The large interface area of the nanoparticles accessible for such interactions is the key for spectacular enhancement in mechanical properties. We shall discuss first the change in glass transition temperature and the dynamic mechanical properties of both covalently and noncovalently polymer functionalized graphene obtained from dynamic mechanical analyser (DMA), then we shall discuss the mechanical properties (stress–strain) for different composite systems.

  Polymer Chemistry Series No. 35 Polymer Functionalized Graphene By Arun Kumar Nandi © Arun Kumar Nandi 2021 Published by the Royal Society of Chemistry, www.rsc.org

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6.2  Dynamic Mechanical Properties The dynamic mechanical properties of a material is a measure of the viscoelastic properties manifested when an oscillatory force is applied to the material. The three important dynamic mechanical properties of the materials are: storage modulus (G′) which relates the ability of the material to store energy, the loss modulus (G″) which relates the ability to lose the energy, and the damping property (tan δ) which is a ratio of G′/G″ observed under application of an oscillatory force to the material. These properties can be measured from the instrument dynamic mechanical analyser where samples are introduced in film form and the property variation on frequency sweep or temperature sweep condition are usually studied.

6.2.1  Covalently Functionalized Graphene Nanocomposites 6.2.1.1  GO-­g PMMA/PVDF Nanocomposite Layek et al.6 used covalently functionalized poly(methyl methacrylate) with rGO(MG) and its composites with PVDF at different filler concentration and the storage modulus, loss modulus and tan δ values of the nanocomposites were measured from the sample film at the tension mode. The samples were heated from 100 to 130 °C at the heating rate of 10 °C min−1. The G′, G″ and tan δ values were measured at a constant frequency of 1 Hz with a static force of 0.02 N (Figure 6.1). From the figure it is apparent that the storage modulus (G′) of the composites is much higher than that of pure PVDF at all temperatures. In the G′ versus T plot (Figure 6.1a) there is a break at about −50 °C, indicating a phase transition (glass transition) in the system. In the loss modulus plot (Figure 6.1b) this glass transition is noticed more distinctly and the transition temperature (Tg) may be computed from the peak temperature. It is apparent from the figure and Table 6.1 that Tg data in the composite has increased by ∼4 °C only; however, if one computes Tg data from the tan δ plot (Figure 6.1c) the maximum increase in Tg is 21 °C. Also the Tg values determined by the two methods differ by 20–37 °C because of two different techniques used for Tg measurement; the former is related to the dissipation of energy as heat (loss modulus) and the latter is related to the reduction of vibration of material, i.e. damping (tan δ). For the former graphene acts as a barrier to heat flow, and thus does not allow dissipation of heat significantly, resulting in a small decrease in Tg. On the other hand, the increased stiffness of graphene composites causes variation of damping to a large extent resulting a large variation in Tg in the composites with MG concentration from the tan δ plot. It is apparent from the Figure 6.1a and b that the storage modulus and loss modulus values are higher in the composites than in pure PVDF, suggesting that the cause of G′ and G″ increase in the MG composites is the same, i.e. due to the reinforcement by MG for its high aspect ratio and good interfacial interaction in the PVDF matrix. To obtain a distinct picture on the increase in G′ in the MG composites from that of pure PVDF, the data shown in Table 6.1 indicate that the percentage increase in storage modulus is maximum

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Figure 6.1  Mechanical  property–temperature plots of different PVDF-­MG nanocomposites: (a) storage modulus, (b) loss modulus and (c) tan δ. Reproduced from ref. 6 with permission from Elsevier, Copyright 2010.

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Table 6.1  Summary  of glass transition temperature (Tg) and storage modulus (G′) values of MG-­PVDF nanocomposites measured by DMA (ref. 6).

G′ G′ G′ (MPa) G′ (MPa) (MPa) at % at −500 % (MPa) % at 500 % Tg Tg Increase at 0 °C Increase °C Increase Sample (°C)a (°C)b −1000 °C Increase °C PVDF MG0.5 MG0.75 MG1 MG3 MG5

−52 −51 −51 −51 −47 −50

−32 −26 −27 −16 −12 −24

4481 6374 7257 7664 8603 8699

42 62 71 92 94

2679 3827 4852 4984 5126 5527

42 81 86 91 106

1068 1458 2395 2240 1927 2176

36 124 109 80 103

653.3 829.9 1453 1319 1011 1129

25 120 100 53 71

a

 easured from loss modulus plots. M Measured from tan δ plots.

b

in the temperature range at the viscoelastic region (−50 to 0 °C) where the movement of the polymer chain segments are relatively free causing reinforcement effect of graphene sheets very large. The highest increase in G′ is 124% observed for the composite containing MG content of 0.75% (w/w) at 0 °C. At this low concentration the uniform dispersion and unidirectional distribution of graphene sheets is responsible for the highest reinforcement. The increase in storage modulus reported for PVDF–clay nanocomposite shows 100%,7,8 for PVDF–Ag nanocomposite the highest increase is 67%,9 and that in a multiwalled carbon nanotube is 120%. Hence, the graphene– PVDF nanocomposite shows more efficiency in reinforcing PVDF from that of nanofillers. So functionalized graphene is a low cost and a highly efficient biocompatible reinforcement filler for PVDF than other nanofillers.

6.2.1.2  GO-­g-­polybenzimidazole/Epoxy Zhang et al.10 tuned the interface of graphene nanoplatelets (fGnPs) by the covalent grafting of polybenzimidazole and made graphene/epoxy composites by the solution casting method. The functionalization of the flakes enabled better dispersion than that of the unfunctionalized flakes and as a consequence higher reinforcement occurs in the composite. The storage modulus versus temperature plot for the samples is presented in Figure 6.2a, and it is apparent that only 0.5 wt% of fGnPs cause about 50% increase in the storage modulus at 30 °C.

6.2.1.3  GO-­g-­polybenzimidazole/PVDF Nanocomposite Maity et al.11 recently used an in situ synthetic technique for the functionalization of polybenzimidazole on graphene oxide (GBI) and produced nanocomposites with poly(vinylidene fluoride) (PVDF) by solution casting technique. GBI exhibits good dispersion in PVDF due to strong interaction of polybenzimidazole with PVDF as evident from FTIR studies. A gradual increase of GBI in the composite increases its piezoelectric β-­polymorph formation with

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Figure 6.2  (a)  DMA curves of fGnPs/epoxy composites. (b) Storage modulus versus temperature plot of PVDF and GBF films. (a) Reproduced from ref. 10 with permission from John Wiley & Sons, Copyright 2018 Wiley Periodicals, Inc and (b) Reproduced from ref. 11 with permission from Elsevier, Copyright 2014.

a maximum of 73% for 10 wt% GBI in PVDF (GBF10) which also exhibits the highest thermal stability. Figure 6.2b shows the variation of storage modulus with temperature for the GBF composite samples. With increasing amount of GBI in PVDF the G′ value increases and shows a maximum value of 1148 MPa for 10 wt% GBI at 30 °C. Thus about 500% increase in storage modulus is noticed in the PVDF composite. The reinforcement of G′ is mainly because of the increased dispersion of filler having high aspect ratio and high interfacial interaction of PBI on graphene surface with the PVDF matrix as evident from FTIR spectra. The GBF10 composite is the optimum composition when the dispersion is uniform with good interfacial interaction On further increase of GBI in PVDF (for GBF15) agglomeration among the GBI sheets

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prevent homogeneous dispersion that decreases storage modulus showing lower reinforcing property. In the case of PVDF the reinforcing property is higher compared to that of epoxy probably due to the different nature of the two polymers and also due to different aspect ratio and composition of the two fillers used by different groups of workers.

6.2.1.4  Ionic Liquid Integrated Graphene/PVDF Composite Maity et al.12 covalently linked imidazolium ionic liquid (IL) with the GO surface to produce GO-­IL which was mixed with PVDF by solution blending, producing PGL composites. The storage modulus (G′), and tan δ plots for pure PVDF and PGL composites with temperature are presented in Figure 6.3a and b. It is

Figure 6.3  Dynamic  mechanical property with variation of temperature plots of PVDF and the PGL composite films. (a) Storage modulus and (b) tan δ. Reproduced from ref. 12 with permission from Elsevier, Copyright 2015.

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clear from Figure 6.3a that, with increase in temperature, G′ of PVDF and PGL composites gradually decreases in a nonlinear fashion and G′ of PGL is much higher than that of PVDF. The G′ values of the composites show a significant increase from that of PVDF and it shows a maximum increased value in the viscoelastic region (0–50 °C) for all composites with a maximum increase for the PGL3 sample (73%). This is due to the reinforcing effect of graphene in the PVDF matrix for its high aspect ratio with uniform dispersion and better interfacial interaction causing easier load transfer, thus increasing the storage modulus in the PVDF matrix.13 In the tan δ versus temperature plot (Figure 6.3b), PVDF shows three peaks at −44.9, 21.5 and 100.2 °C corresponding to three transition temperatures (Tg, Tr(I) and Tr(II)), where Tg is the glass transition temperature, Tr(I) is the relaxational transition of the PVDF chains present at the amorphous–crystalline interface, experiencing a partial constrained state for movement, and the Tr(II) represents the relaxation temperature of PVDF chains at the crystalline zone where the segments are in a constrained state.14–16 This is the probable reason for Tg < Tr(I) < Tr(II). It is interesting to note that the Tg and all other transition temperatures of PVDF increases with increase in GO-­IL concentration in the composite and may be attributed to the increase in compactness of the respective zones of PVDF due to the strong attractive force between the GO-­IL nanosheets, which are well dispersed in the composite.17 Also the formation of the β-­crystal phase of PVDF in the presence of GO-­IL may be a possible reason because the β-­phase has more compactness than the α-­phase, so it requires a higher temperature for the relaxation process to occur.

6.2.2  N  oncovalently Functionalized Graphene Nanocomposites 6.2.2.1  P  olythiophene-­graft-­poly(Methyl Methacrylate) RGO/ PVDF Composites In a novel attempt, Maity et al.18 noncovalently functionalized polythiophene-­ graft-­poly(methyl methacrylate) (PT-­g-­PMMA) with RGO via a π-­stacking interaction, producing PFG which was used to make composites with PVDF, designated as PRPx where x represents the weight percentage of PFG content. The RGOs were found to be distributed homogenously in the PVDF matrix. To study the influence of PFG on the dynamic mechanical properties of PVDF the storage modulus (G′) and tan δ plots of the composites are presented in Figure 6.4 for varying temperature at a constant frequency of 1 Hz. From the figure it is apparent that the G′ of PRP composites gradually decreases in a nonlinear way with rise of temperature and at a fixed temperature G′ of the composites gradually increases with increase in PFG concentration showing a maximum of 60% increase at 25 °C for the PRP5 sample (Table 6.2). The increase in G′ is attributed to the reinforcing effect of graphene for its sheet-­like structure having high aspect ratio, uniform dispersion and strong interfacial adhesion between the filler and PVDF matrix. The strong interaction between the >C=O group of the PMMA unit of graft co-­polymer

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with the >CF2 dipole at the interfacial region of PVDF facilitates easier energy transfer, causing enhancement of mechanical properties.17 The tan δ versus temperature plot for neat PVDF and PRP composites are also presented in Figure 6.4 where each curve exhibits two prominent peaks and one hump at the intermediate temperature region. The lower temperature peak of the tan δ curve of PVDF at −46.2 °C is for the Tg of PVDF, and on addition of PFG, Tg peaks shift to higher temperatures from −46.2 to −42.7 °C for the PRP5 sample.17,19 The Tg is related to the response to the confinement of segmental motion, which increases with PFG loading due to the strong cohesive force between the PFG sheets for its large surface area making a more compact state of the polymer thus decreasing the free volume of the polymer necessary for segmental motion. The increments in the Tg values with increasing PFG content illustrates the above confinement effect.6 The tan δ versus T plot of Figure 6.4 exhibits one hump at 21.4 °C and another peak at 88.2 °C for PVDF identifying the transition temperatures at the crystal–amorphous interface (Tr(I)) and the crystalline–crystalline relaxation zone (Tr(II)), respectively. All the transition temperatures of the PRP composites gradually increases over the PVDF due to the surface force of PFGs dispersed in the PVDF matrix (Table 6.2). The increase in Tr(I) may be attributed to the hindered relaxation process at the crystalline–amorphous zone of PVDF and to some extent for the β-­polymorphic structure produced in the composite.6,17,20 The relaxation temperature (Tr(II)) of PVDF is also ascribed to the segmental motion of PVDF polymer chain within the crystalline zone and with increase in PFG content it shifts to higher temperature. Thus, the combined effect of interfacial compatibility arising from the noncovalent functionalization of

Figure 6.4  Storage  modulus and tan δ versus temperature plots of PVDF and the

PRP composites Reproduced from ref. 18 with permission from Elsevier, Copyright 2016.

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Table 6.2  Glass  transition (Tg), crystal–amorphous (Tr(I)) and crystal–crystal (Tr(II))

transition temperatures and storage modulus (G′) values (from DMA) of PVDF and their PRP composites (ref. 18).

Samples

Tg (°C)

Tr(I) (°C)

Tr(II) (°C)

G′ (MPa) at 25 °C

% Increase in G′

PVDF PRP0.l PRP0.3 PRP0.5 PRPl PRP3 PRPS

−46.2 −43.5 −43.8 −44.2 −43.0 −43.5 −42.7

21.4 24.4 28.2 31.1 30.1 31.6 33.2

88.2 90.3 97.0 104.1 103.1 103.0 98.0

1468 1945 1960 1907 2005 1994 2350

— 32.5 33.5 30.0 36.6 36.0 60

graphene by PT-­g-­PMMA and homogeneous distribution of PFG in the PVDF matrix increases the reinforcing property as well as cause increase of different transition temperatures.20

6.2.2.2  Functionalized GO/Epoxy Nanocomposite Ferreira et al.21 noncovalently functionalized GO with hexamethylenediamine and studied the mechanical properties of epoxy-­based nanocomposites using DMA to assess the information on the viscoelastic properties of the composites. It was found that only 20% increase in storage modulus in the unfunctionalized GO/epoxy nanocomposite is noticed. However, in the amino functionalized GO/epoxy nanocomposite about 30% increase in storage modulus is observed due to the reinforcing effect of graphene on the matrix. The addition of graphene (modified or not) to the epoxy matrix leads to a phase formation at the interface between the filler–matrix, which contributes to dissipating energy from external stresses by frictions between particle–particle and particle–polymer at the interface. The tan δ value of GO/ epoxy nanocomposites decreased slightly from that of neat epoxy and for the amino functionalized GO loaded epoxy system the lowest value for tan δ was noted. It probably occurs because the addition of fillers affects the damping behaviour of the polymer owing to stress concentrations.22 Also, a good adhesion between graphene and epoxy resin limits the mobility of the polymer chains, reducing the values of tan δ.23 In this system Tg of the composites shifts to 6–8 °C lower temperature than that of pure epoxy. The reduction in Tg is attributed to the enlarged free volume between the epoxy chains for the introduction of the nanofiller into the epoxy polymer.24 Fang et al.25 made supermolecular aggregates of piperazine (PiP) and phytic acid (PA) placed onto the GO surface in water to fabricate functionalized GO (PPGO) to make composites with epoxy resin (EP). Figure 6.5 exhibits the influence of temperature on storage modulus (E) values and loss factor (tan δ) of EP, EP/GO3 and EP/PPGO3 where the number indicates the wt% of the filler. In Table 6.3 the storage modulus at 30 °C (glassy state), at Tg + 30 °C (rubbery state) and glass transition temperature (Tg) are presented. At 30 °C,

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Figure 6.5  Storage  modulus values (E) and loss-­tangent (tan δ) at 1 Hz of cured EP, EP/GO3 and EP/PPGO3. Reproduced from ref. 25 with permission from Elsevier, Copyright 2019.

Table 6.3  Storage  modulus at 30 °C, 30 °C above the glass transition temperature (Tg) and Tg of EP, EP/GO3 and EP/PPGO3 (ref. 25).

EP EP/PPG03 EP/G03

E′ at 30 °C (MPa)

Er (E′ at Tg + 30) (MPa) 3

Tg (°C)

1614 1944 1599

14 51 44

134 142 141

EP/PPGO3 shows an E′ value of 1944 MPa, 20% higher than that of neat EP (1614 MPa) and 27% higher than that of EP/GO3 (1599 MPa). The increased E′ indicates that EP/PPGO3 is more rigid than neat EP and EP/GO3, suggesting that PPGO exhibits a better reinforcing effect than GO in this system. This is due to the stack of GO sheets on self-­assembling with peperazine and phytic acid, when this nanoscale roughness causes an enhanced mechanical interlocking with the polymer chains.26 Relatively, a higher storage modulus (Er) at the rubbery state of EP/PPGO3 and EP/GO3 than that of EP is observed, due to stronger interaction at the interface between graphene and matrix. Meanwhile, the higher Er of EP/PPGO3 than that of EP/GO3 further indicates the better dispersion of PPGO in the EP matrix. In the rubbery state a 250% increase in storage modulus is observed in EP/PPGO3 composite while in EP/ GO3 the increase is about 200%. This shows a higher reinforcing effect of the fillers at the rubbery state than that of glassy state (30 °C) probably due to stronger interfacial interaction at the rubbery state than that of the glassy state. The Tg, calculated from the maximum of tan δ, increases by 7–8 °C from that of neat EP and follows the same trend as in the E′ value, contrary to the earlier graphene/epoxy composite. A probable reason may be due to the different nature of filler and epoxy that causes stronger interaction in the latter than that of the former system.

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6.2.3  Functionalized Graphene/Polystyrene Composites Using a sonochemical method, Wang et al.27 prepared styrene functionalized graphene (FG) by direct exfoliation of graphite flakes in the styrene monomer. This material is dispersed in toluene and produced composites with polystyrene (PS) by solution casting followed by a compression moulding method. The dynamic mechanical properties of neat PS and PS/FGs composites are measured by scanning with temperature at a constant frequency of 1 Hz (Figure 6.6a and b). Figure 6.6a shows the effect of FG concentration on storage modulus and it increases with increasing FG concentration at a constant temperature showing a decline at 1.0 wt% loading of FG. The enhancement in the storage modulus indicates that FG is acting as an effective reinforcing agent in the PS matrix by transferring the load from the polymer. It is noted that a maximum increase in storage modulus occurs with a value of ∼41% corresponding to 0.5 wt% of FGs loading at 40 °C. The increase in storage modulus of the PS/FGs composites is attributed to the uniform dispersion of FG sheets in the PS matrix, together with π-­stacking interaction between the FG and PS matrix. From Figure 6.6b the Tg, measured from the tan δ plot, increases with FG content, indicating that the segmental mobility of the PS chains becomes hindered during glass transition by the presence of FGs. The maximum shift of Tg is 9.2 °C (from 87.7 to 96.9 °C) for 1.0 wt% FGs content. The origin of Tg shifts is attributed to the presence of the so-­called ‘interphase’ polymer, formed due to the interaction of the PS chains with FG surface and percolation of this network of interphase polymer is responsible for the hike in Tg of the bulk composite.28

Figure 6.6  (a)  Storage modulus and (b) tan δ curves of PS/FGs composites. Reproduced from ref. 27 with permission from the Royal Society of Chemistry.

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6.3  Mechanical Properties The mechanical (stress–strain) properties of polymers are very important, particularly for their application in different appliances. These include tensile strength, tensile strain, Young's modulus, toughness, etc. Which are very important for their specific application in both commodity and engineering uses. Due to the high aspect ratio of nanoparticles and fruitful interfacial adhesion with the polymer matrix, stress transfer throughout the whole polymer is possible via nanoparticles. This causes a significant improvement in the mechanical properties and in this respect graphene or functionalized graphene is a more effective nanofiller than conventional fillers used to make polymer composites. Like earlier, here we shall discuss the mechanical properties of polymer nanocomposites produced with polymer functionalized graphene made by covalent and noncovalent paths in the following sections.

6.3.1  Covalently Functionalized Graphene Nanocomposites 6.3.1.1  GO-­g-­PMMA/PVDF Nanocomposite Layek et al.6 made stress–strain curves of PVDF graphene nanocomposites containing PMMA covalently grafted on the GO surface(GO-­g-­PMMA) and these are presented in Figure 6.7a. It is apparent from the figure that for a particular value of strain there is an increase in tensile stress with increasing MG content, indicating an increase in film stiffness for the addition of MG. The elongation at break gradually decreases with increase in graphene content. The Young's modulus, stress at break, strain at break and their percentage increase with graphene content are presented in Table 6.4. The Young's modulus (Ey) of PVDF increases from 1.44 GPa to 6.07 GPa for 5% modified graphene (MG) and the percentage increase is really high (321%) at this filler concentration. The stress at break also increases very dramatically from 39 to 100 MPa for the 5% (w/w) MG sample with a percentage increase of 157%. The tensile stress and tensile strain values plotted with MG concentration are presented in Figure 6.7b, which shows the exponential increase of tensile stress at break and exponential decrease in tensile strain with respect to MG concentration. This figure also indicates at 3% (w/w) MG concentration the levelling value of tensile strength and tensile strain, indicating a rigidity percolation threshold at this composition. The toughness values presented in Table 6.4 shows a gradual decrease with increase in MG content. In Figure 6.8 the Young's modulus (Ey), measured from the initial slope of the stress–strain curve, is plotted with MG content and it initially increases with MG concentration slowly, then increases rapidly at the concentration range 0.5–1%, and finally the rate of increase is slow resembling an autocatalytic (cooperative) nature (inset). This very large increase in Ey in this composite may be attributable to the uniform dispersion of MG sheets in the composite due to blending. The Halpin–Tsai equation29 is used to calculate the effect of volume fraction of filler on the Young's modulus of the

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Figure 6.7  (a)  Stress–strain curves of the pure PVDF and PVDF-­MG nanocomposite

films with different contents of modified graphene sheets; (b) Mechanical properties of PVDF-­MG nanocomposite films with different MG contents: tensile strength (left) and elongation at break (right) versus graphene loading. Reproduced from ref. 6 with permission from Elsevier, Copyright 2010.

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Table 6.4  Mechanical  properties (stress, strain, modulus, toughness) of PVDF and PVDF-­MG nanocomposites at 300 °C (ref. 6).

% % % Stress Increase Young's Increase Decrease Strain at at break of stress modulus of Young's Toughness of Sample break (MPa) at break (GPa) modulus (MPa) toughness PVDF 16.5 MG0.5 10.5 MG0.75 8.6 MG1 7.5 MG3 5.4 MG5 4.6

39.2 67.3 77.3 82.9 97.5 100.7

— 71 97 111.3 148.4 156.5

1.44 2.25 3.19 4.04 5.26 6.07

— 77 121 180 265 321

6278 5301 4914 5028 4210 3898

— 15 22 20 33 38

Figure 6.8  Young's  modulus versus vol% MG of the PVDF-­MG nanocomposite

films and Halpine–Tsai theoretical (random and unidirectional) plots. Reproduced from ref. 6 with permission from Elsevier, Copyright 2010.

constituents in a simple way and attempts are now being made to explain the Young's modulus data of graphene composites with the help of the equation.30,31 Considering the random distribution of the MG in the composite, the modified form of the Halpin–Tsai equation may be written as:32–35   

 3 1  (2 LG /3TG ) L VG 5 1  2T VG   E    EP 1   L VG 8 1  T VG     8

where

L 

E

E G

G

EP   1

EP    2 LG 3TG 

(6.1)

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178

and

T 

E E

G

G

EP   1

EP   2

where Eγ, EG and EP are the Young's modulus of the composite, graphene and polymer, respectively. LG, TG and VG are the length, thickness and volume fraction of graphene units present in the composites. Thickness values of MG are obtained from the AFM data in the composite. The above equation is developed for fibre/tubular morphology of nanofiller and here the effective length of the graphene sheet is obtained from measured average length, breadth and thickness of MG sheet from AFM data of the MG3 composite. The average length, breadth and thickness have values of 50, 15 and 1.5 nm, respectively. The graphene platelet is regarded as being made of ten rectangular parallelepipeds (15/1.5 = 10) contributing to a total length (LG) of 50 × 10 = 500 nm and thickness (TG) of 1.5 nm.6 VG is calculated from the weights of graphene and PVDF in the composites taking density (ρ) of PVDF as 1.92 g cm−3 and ρgraphene = 0.80 gm cm−3. For unidirectional (parallel) orientation of graphene nanosheets in the polymer films the modified Halpin–Tsai equation is represented by the equation:   



  

 1  (2 LG /3TG ) L VG  EYn   . 1   L VG  

(6.2)

From Figure 6.8 it is evident that the experimental Ey values lie between the two theoretical curves drawn assuming random and unidirectional distribution. An analysis of the experimental Ey data suggest that at low concentration of graphene it is near to the unidirectional (parallel) orientation but at high graphene concentration (>1% v/v) it approaches random orientation. A possible explanation for the parallel orientation at low MG content is due to the directional nature of the specific interaction between PVDF and PMMA at lower MG concentration, causing unidirectional (parallel to the film) orientation of the MG sheets. As the MG concentration increases, the interaction of PVDF on a large number of MG sheets loses its directional character and becomes isotropic. This results in a random distribution of MG sheets at higher MG concentration in the composites. The experimental Ey curve suggests that the transition from parallel to random orientation of graphene sheets is cooperative in nature, signifying that the directional character of interaction changes drastically by a small increment in MG concentration in the composite.

6.3.1.2  f-­(PVA)GO/PVA Composite Cano et al.36 covalently functionalized GO with poly(vinyl alcohol) (PVA) by esterification of carboxylic groups of GO with hydroxyl groups of PVA forming functionalized f-­(PVA)GO as paper-­like material. Papers prepared from

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f-­(PVA)GO exhibit significant enhancement of mechanical properties compared to those prepared with GO or with simple mixtures of GO and PVA. The best performance has been achieved for PVA with molecular weights between 50 and 150 kg mol−1. Enhancement in Young's moduli by 60% and tensile strength by 400% are observed relative to GO-­PVA mixture. The improved mechanical properties are explained from enhanced inter-­flake stress tran­s­fer for the covalently bonded PVA. This functionalized f-­(PVA)GO was used as filler in PVA-­based composites and the representative tensile stress–strain curves for f-­(PVA)GO in PVA composites are presented in Figure 6.9A. It is

Figure 6.9  Mechanical  properties of composites prepared from f-­(PVA)GO (Mw = 50 kg mol−1) dispersed in a PVA matrix (Mw = 78 kg mol−1). (A) Representative stress strain curves. (B) Young's modulus. (C) Ultimate tensile strength. (D) Strain at break. Reproduced from ref. 36 with permission from Elsevier, Copyright 2012.

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evident from these curves that the addition of f-­(PVA)GO has increased both the modulus and tensile strength of these composites. Similar curves for pristine GO–PVA composites do not exhibit any reinforcements. In Figure 6.9B the Young's modulus data for the composite as a function of f-­(PVA)-­GO volume fraction is shown. The modulus has increased from 2.5 GPa for PVA to 4.5 GPa for the 0.24 vol% filler linearly, beyond which the modulus falls off. The rate of linear increase is dY/dVf = 610 GPa. In a similar fashion, the strength has increased from 90 MPa for PVA to 140 MPa for the 0.24 vol% filler before falling off, and the slope of the linear region is 22 ± 2 GPa. These significant increases are higher than the best published results for polymers filled with either GO or pristine graphene.37–39 In Figure 6.9C is the strain at break as a function of filler content and no significant impact on the material ductility by the f-­(PVA)-­GO filler is noticed. The data for the others with pristine GO composites are also shown by an open symbol in Figure 6.9B–D, which illustrates lack of reinforcement in unfunctionalized GO–PVA composites.

6.3.1.3  H  yperbranched Polyamide Functionalized GO–Epoxy Nanocomposites Li et al.40 used amine-­terminated hyperbranched polyamide (HBPA-­NH2) covalently grafted on GO (GO-­HBPA) to improve the mechanical properties of the GO/epoxy and the GO-­HBPA/epoxy nanocomposites. Figure 6.10a–d show the tensile strength, elongation at break, tensile modulus, brittleness, flexural strength, and flexural modulus of the nanocomposites. The tensile strength along with the tensile modulus of the GO–epoxy nanocomposites and the GO-­HBPA-­epoxy nanocomposites indicate the same trend (Figure 6.10a), i.e. they first increase and then decrease as the nanofiller loading increases. The latter decrease is due to the aggregation of GO or GO-­HBPA in the epoxy matrix at high nanofiller concentration. At 0.15% GO-­HBPA concentration, the tensile strength of the GOHBPA–epoxy nanocomposites adopts the maximum value of 69.9 MPa, which is 42.3% greater than the value of neat epoxy. For the GO–epoxy nanocomposites, the highest tensile strength is 56.8 MPa at GO loadings of 0.10% and the tensile strength of the GO-­HBPA-­modified nanocomposites is higher at the same nanofiller loadings, suggesting that the strength enhancement effect of the GO-­HBPA is better than that of GO. These results can be attributed to the better dispersity of GO-­HBPA than GO. The elongation at break and the brittleness of the nanocomposites are presented in Figure 6.10b and c, where the elongation at the break of the GO–epoxy nanocomposites and the GO-­HBPA–epoxy nanocomposites increases first and then decreases as the nanofiller loading is increased, similar to the trend of tensile strength. As for the GO-­modified nanocomposites, the elongation at break reached the maximum of 6.5% at 0.10% GO loading and it is 109% higher than that of neat epoxy. When 0.15% GO-­HBPA is added, the elongation at break has reached a maximum value of 16.4%, which is 429% higher from that of neat epoxy and 152% higher from

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that of the GO-­modified nanocomposites. The much higher elongation at break of the GO-­HBPA-­modified nanocomposites at the same nanofiller loadings reveals that the toughening effect of the GO-­HBPA is better than that of the GO. Figure 6.10c shows that GOHBPA at 0.15% loading has the lowest brittleness in the nanocomposites, which is similar to the elongation at break results. So, the GO-­HBPA-­0.15% possesses the highest mechanical strength and the best toughness among the nanocomposites. Here, the enhancement of tensile strength and elongation at break of the present epoxy nanocomposites are higher than most previous works41,42 indicating that GO-­HBPA prepared in this work plays a significant role in the improvement of fracture toughness and mechanical strength of epoxy nanocomposites. The flexural strength and flexural modulus of the neat epoxy and graphene/epoxy nanocomposites are presented in Figure 6.10d, which shows significant improvement in the flexural strength and flexural modulus of the nanocomposites with addition of GO-­HBPA. In comparison with the neat epoxy, the flexural strength of the GO-­0.06% has increased by 1.6% (113.8 to 115.7 MPa), and the flexural modulus has enhanced by 4.0% (2.96 to 3.08 GPa). Remarkably, at 0.15% GOHBPA loading, the flexural strength of the epoxy nanocomposites

Figure 6.10  Mechanical  properties of the neat epoxy and graphene/epoxy nanocom-

posites: tensile strength and tensile modulus (a), elongation at break (b), brittleness (c), flexural strength and flexural modulus (d). Reproduced from ref. 40 with permission from Elsevier, Copyright 2019.

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has increased by 41.3% (113.8 to 160.8 MPa), and the flexural modulus has increased by 97.3% (2.96 to 5.84 GPa.). So the enhancement value for GO-­ HBPA nanofiller is much higher from that of GO nanofiller and it is explained from the well dispersion of the GO-­HBPA in the epoxy matrix and the good interfacial interactions between the GO-­HBPA and the matrix. Wan et al.43 used diglycidyl ether of bisphenol-­A functionalized GO (DGEBA–f–GO) sheets to produce epoxy composites. The functionalization of the DGEBA layer effectively improves the compatibility and dispersion of GO sheets in epoxy matrix. The tensile stress–strain experiment shows that the DGEBA-­f-­GO/epoxy composites exhibits higher tensile modulus and strength than either neat epoxy or pristine GO/epoxy composites. For an epoxy composite with 0.25 wt% DGEBA-­f-­GO, the tensile modulus and strength increased from 3.15 to 3.56 GPa (∼13%) and 53 to 93 MPa (∼75%), respectively, from the neat epoxy resin.

6.3.1.4  Ionic Liquid Integrated Graphene/PVDF Composite A covalently linked imidazolium ionic liquid (IL) from the GO surface (GO-­ IL) was used by Maity et al.12 to mix with PVDF by solution blending process to produce PGL composites. The enhancement of mechanical properties of the composite films at different compositions of GO-­IL was evaluated from the stress–strain measurement (Figure 6.11a). From the stress–strain curves it is evident that the mechanical properties of the PGL composites exhibit a significant change from that of a pure PVDF matrix. The Young's modulus, tensile strength and elongation at break of all composites with respect to volume fraction of GO-­IL measured from the stress–strain plots of four different measurements and the average values are presented in Figure 6.11b–d along with their standard deviations. The filler concentration is converted from mass fraction ‘w’ (wt%) to volume fraction ‘V’ of GO-­IL using the density values of PVDF and graphene nanosheets, as 1.92 g cm−3 and 2.2 g cm−3, respectively.6 With increasing volume fraction of GO-­IL the Young's modulus of the composite gradually increases; however, the elongation at break gradually decreases. For 2.63% (v/v) addition of GO-­IL, the Young's modulus (Ey) of PVDF increases from 1.15 to 4.98 GPa yielding a 333% increase, which is quite high. The tensile strength of the samples also significantly increases from 9.96 to 72.54 MPa at this composition with an increase in 628% and it can be attributed to the formation of three-­dimensional network structure present in the PGL composites. The elongation at break of PGL films gradually decreases from that of PVDF with the increase in GO-­IL concentration for the large aspect ratio and good interfacial interaction of graphene nanosheets causing the composite to be stiffer.17,44,45 Here the Halpin–Tsai equation is used to theoretically understand the Young's modulus (Ey) for the random or unidirectional distribution of filler in the composites. It is observed that at low concentration, the graphene nanosheets obey the random distribution as they are well-­dispersed for the isotropic nature of supramolecular interaction between the components but at higher concentration it is closer to the parallel (unidirectional) orientation

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Figure 6.11  Mechanical  properties of PGL composite films: (a) stress–strain

curves, (b) Young's modulus versus loading of GO-­IL (vol%), (c) tensile strength and (d) elongation at break. Reproduced from ref. 12 with permission from Elsevier, Copyright 2015.

to the surface of the composite films. This is because at higher loading of GO-­ IL they lose their isotropic nature and on applying external tensile loads it is successfully transmitted from the fillers across the PVDF– graphene interface causing the unidirectional distribution of GO-­IL.44–47 From the above results it is inferred that the random orientation of filler changes to parallel orientation by increase of GO-­IL in the PGL composites.

6.3.1.5  P  oly (2-­Hydroxyethyl Methacrylate) Functionalized Graphene (PHEMA-­G)/Poly(p-­phenylene Benzobisoxazole) (PBO) Composite Hu et al.48 used the ATRP method to attach poly(2-­hydroxyethyl methacrylate) (PHEMA) from exfoliated graphene surface where ATRP initiators are introduced to pristine graphene via a one-­step carbene reaction and it is further polymerized by ATRP to produce different graphene/PBO composites.

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Grafting of PHEMA improved the dispersity, interfacial adhesion and reactivity of graphene in the PBO matrix. It was blended with PBO via in situ polymerization of p-­phenylene benzobisoxazole. The hydroxyl groups of HEMA help to form covalent bond between PBO fibre and nanofiller and graphene/PBO composite fibres are fabricated with the help of a dry-­jet wet spinning process. To understand the enhancement of mechanical properties of PBO on adding PHEMA-­G the tensile strength is compared with that of pristine graphene/PBO with different content of PHEMA-­G, which is presented in Figure 6.12a. The tensile strength of composite fibres increases

Figure 6.12  (a)  Tensile strength changes of PBO composite fibres with increasing

graphene content. (b) Young's modulus of G/PBO composite fibres and the theoretical results from the Halpin–Tsai equation. Reproduced from ref. 48 with permission from American Chemical Society, Copyright 2017.

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with graphene loading in the matrix and in comparison with the pristine graphene composite, the PHEMA-­G exhibits more prominent enhancement of tensile strength. It is interesting to note that the tensile strength of the composite fibre is enhanced by 51.2% (from 2.13 to 3.22 GPa) at 1.0 wt% loading of PHEMA-­G. This increment is much higher compared to that of 1.0 wt% modified carbon nanotubes49 (34.5% increase) and pristine graphene50 (11.9% increase), respectively. This outstanding improvement in the tensile strength of PHEMA-­G is attributed to: (i) better dispersion in the PBO matrix, (ii) formation of covalent links between PHEMA-­G and PBO chains and (iii) the nondestructive modification strategy of graphene. The Young's modulus for the composite fibres also exhibits a similar increase as the tensile strength (Figure 6.12b). For the composite fibre with 1.0 wt% of PHEMA-­G, the Young's modulus increases from 104 to 139 GPa, showing 33.7% increase. With such a low PHEMA-­G content this significant increase is very noteworthy; however, on further increasing PHEMA-­G loading, at 1.5 wt% a sluggish increase rate of both Young's modulus and tensile strength is noted. At the high PHEMA-­G loading the fillers restack, causing agglomeration, thus saturating the reinforcing efficiency. Also, the crystallization behaviour can affect the tensile strength and Young's modulus of the PBO fibres. The decreased crystal size may also be a cause for the sluggish increase rate. In order to understand the PHEMA-­G dispersion and orientation in the PBO fibres, Halpin–Tsai equations are applied.50,51 For this purpose the Young's modulus of the graphene and PBO are taken as 1 TPa and 104 GPa, respectively, thickness and average length of graphene sheets are 2.5 nm and 1 µm according to the AFM and TEM analysis; the density of graphene and PBO are taken as 1.60 and 1.56 g cm−3, respectively. Figure 6.12b exhibits that the Young's modulus of composite fibres is much higher than that of theoretical predictions indicating graphene is unidirectionally aligned along the fibre axis. These results therefore suggest that PHEMA-­G is highly aligned in the PBO matrix produced by the dry-­jet wet spinning process.

6.3.2  N  oncovalently Functionalized Graphene Nanocomposites 6.3.2.1  P  olythiophene-­graft-­poly(Methyl Methacrylate) RGO/ PVDF Composites Maity et al.18 noncovalently functionalized polythiophene-­graft-­poly(methyl methacrylate) (PT-­g-­PMMA) with rGO producing PFG to make composites with PVDF, designated as PRPx where x represents weight percentage of PFG content. To understand the effect of PFG on the mechanical performance of PVDF/graphene composites, the typical stress–strain plots for different PFG content are shown in Figure 6.13a–c. Figure 6.13a indicates that both tensile strength and tensile strain increase in the PRP composites from those of neat PVDF, and in Figure 6.13b these values are compared for different PFG content in the composites. Initially the tensile strain increases with increase of

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Figure 6.13  Mechanical  properties of PRP composite films: (a) stress–strain curves,

(b) variation of tensile strength and tensile strain of the PRP samples (wt%) and (c) Young's modulus versus vol% of PFG in the PRP composite samples. Theoretical data using the Halpine–Tsai parallel and random models is used to correlate with the experimental data. Reproduced from ref. 18 with permission from Elsevier, Copyright 2016.

PFG content and exhibits a maximum at ∼1 wt% PFG content, and on further addition of PFG it decreases in the PRP composites. A probable reason is that at low PFG concentration the random distribution of the filler induces the PVDF chains to uncoil more; however, at higher PFG concentration some aggregation and change of filler distribution from random to oriented structure (discussed below) results in the formation of stiffer film. Likewise, the tensile strength initially shows a sharp increase showing a maximum value of 355% for the PFP5 sample. It is to be noted that the PRP1 sample exhibits increment of tensile strength 317% and tensile strain 302% from PVDF, respectively. This large enhancement of both tensile strain and tensile strength indicates the improved ductility of PVDF at PRP1 composite. The Young's modulus (Ey) of PVDF increases from 1.15 to 7.89 GPa for 5 wt% PFG

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in PVDF showing a maximum increase of ∼585% for the PRP5 composite. This is because the stress required per unit strain is higher for the more compact nature of the sample and PFG acts as a primary load-­bearing component of the PVDF composite acting as effective stress transfer agent in the composite increasing the mechanical properties significantly.6 Thus, PT-­g-­PMMA stabilized graphene exhibits concomitant enhancement of Young's modulus and tensile strength of PVDF, therefore it can act as a structural material for domestic and engineering applications. The modified Halpin–Tsai equation is widely used for predicting Young's modulus (Ey) of composites for unidirectional or random distribution of the functional fillers for two different orientation,44,52 e.g. parallel to the surface of the PVDF film or randomly distributed within the PVDF matrix. The Young's modulus of the PRP composite films are calculated theoretically with the Halpin–Tsai equation taking Young's modulus of the graphene sheet and PVDF as 1 TPa and 1.15 GPa,6 respectively The measured average length (LG) and thickness (TG) of the graphene sheet are 700 nm and 0.9 nm, respectively and VG is calculated according to the method described above. In Figure 6.13c the theoretical and experimental data are plotted with volume fraction of PFG and at low PFG concentration (CF2 dipoles of PVDF is not sufficiently rapid to reach the equilibrium with the applied field at higher frequencies, causing a sharp decrease in the dielectric permittivity with increasing frequency.11 In the inset of Figure 8.1 a plot of dielectric permittivity at 102 Hz is plotted with the GO-­IL content in the composites and a sharp rise at the GO-­IL concentration of 0.1% (w/w) is observed. It is important to note that this value of percolation threshold is much lower than that for the poly(vinylidene fluoride-­co-­hexafluoropropylene)/rGO composite (0.8%).7 So lowering of the percolation threshold probably occurs for the supramolecular organization of PVDF/GO-­IL composite producing a fibrillar

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network morphology producing large density microcapacitors. It is amazing that the present dielectric threshold value matches well with the percolation threshold value obtained from dc conductivity presented in Chapter 7 and this matching suggests that for both the conductivity and dielectric permittivity the same reason exists, i.e. the presence of a three dimensional network in the PGL composites.

8.2.1.2 Polyaniline Functionalized Graphene/PVDF Composites Graphene based conducting materials are prime candidates for energy applications due to augmentation of intriguing properties that triggers a great area of academic research and industrial applications.12 So, to increase the dielectric constant of PVDF more significantly, Maity et al.13 fabricated hierarchical nanostructured PGP composites of PVDF with polyaniline covalently grafted with graphene oxide (G-­graft-­PANI) and during the grafting process it is converted to rGO. The HR-­TEM images exhibit homogeneous dispersion of rGO in PVDF and FE-­SEM images show fibrillar morphology generated from supramolecular assembly of G-­graft-­PANI with PVDF chains resulting hierarchical nanostructure of PGP composites. The dipolar interaction between the filler and PVDF, and the high surface area of G-­graft-­PANI transforms the nonpolar α-­phase PVDF to 91% polar β-­phase at 5 wt% filler content. The variation of dielectric permittivity (ε′) with frequency are shown in Figure 8.2a of the PVDF composites for varying G-­graft-­PANI concentration measured at 30 °C. PVDF exhibits a dielectric permittivity value of 10 which is almost frequency independent at the measured frequency range of 102–106 Hz. As the filler weight percentage increases the dielectric permittivity of PGP composites increases gradually and reaches a value of 264 for 5 wt% filler at 102 Hz showing an enormous (2500%) increase. Here, G-­graft-­PANI nano-­fillers are comprised of two conducting components, e.g. graphene and PANI. This enormous increase in dielectric permittivity of the PVDF composites with the G-­graft-­PANI filler is attributed to the interfacial polarization occurring at the gradient interface between the G-­graft-­PANI layer and the PVDF surface region. It is necessary to mention that the dielectric permittivity of neat G-­graft-­PANI is very high (3 × 105)13 and on its introduction into the PVDF matrix to an extent of 5 wt% results in an enormous increase in ε′ making it suitable for different applications. A comparison of PVDF composites with noncovalently functionalized PANI on GO surfaces exhibit a small increase of ε′ from 7 of pristine PVDF to ∼10 at 5 (vol%) nanofillers14 and Li et al. reported a maximum ε′ value of ∼125 at 102 Hz for the PVDF-­PANI system for 13 (vol%) PANI loading.15 However, the present PVDF/G-­graft-­PANI nanocomposite exhibits a maximum ε′ value of 264 with a filler loading of 5 wt%, making it superior to other reported systems. This enormous increase in dielectric constant is attributed to the Maxwell–Wagner–Sillars interfacial polarization of different conductive layers and specific interactions between PVDF and the G-­graft-­PANI surface.16 The increase in ε′ is also dependent on the high

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Figure 8.2  (a)  Variation of dielectric permittivity with frequency, (b) dielectric per-

mittivity and tan δ with different compositions of G-­graft-­PANI in PVDF at 30 °C at 100 Hz. (c) Fitting of dielectric permittivity with eqn (8.1). (d) Variation of tan δ with frequency of PVDF and PGP composites of different composition. Reproduced from ref. 13 with permission from Elsevier, Copyright 2016.

aspect ratio of the graphene sheets and their uniform distribution as well as their orientation in the PVDF matrix. According to this polarization theory, when an electric field is applied to the composite film the charge carriers originate from the external electrode and then migrate in conductive G-­graft-­ PANI layers where they become stored at the interface between conducting G-­graft-­PANI and insulating PVDF matrix for the difference in their relaxation time. This storing of charges at the interface contributes to a large interfacial polarization yielding increased permittivity. The gradual increase in dielectric permittivity with filler concentration may be attributed to the formation of increased number of microcapacitor networks by the conducting G-­graft-­ PANI layers in the PVDF matrix. It is to be noted that, with increase in frequency, there is a decrease in dielectric permittivity and this is ascribed to the change in orientation of dipoles rotating rapidly, with a persistent lag

Chapter 8

238

between the frequency of the oscillating dipoles and that of the applied field. FTIR and XRD results indicate that polar β-­phase formation in the composites of G-­graft-­PANI and PVDF and this polymorphic form is favourable for the augmentation of dielectric properties of the composites.17 In Figure 8.2b the bar diagram indicates the variation of dielectric permittivity and tan δ (dielectric loss) with concentration of G-­graft-­PANI. The dielectric permittivity increases significantly when the filler content approaches the ‘percolation threshold’ at 1 wt% which is attributed to the development of the dielectric G-­graft-­PANI layer. The change in dielectric permittivity (ε′) with weight percentage of the nanofiller may abide by a power law which can be elucidated by using the percolation theory:18   

  

ε' ∝ εm ( fc − f )−s, where fc > f

(8.1)

where εm is the dielectric permittivity of PVDF and f and fc are the wt% of fillers in the composites and at the percolation threshold of the fillers in the matrix, respectively, and s is the critical exponent. The best linear fit of the dielectric data in the log ε′ versus log(  fc − f) plots, is shown in Figure 8.2c with fc = 1 (wt%), obeying the power law showing the exponent (s) value = 1.24. This low percolation threshold suggests a good dispersion of G-­graft-­PANI in the gradient interface of PVDF matrix as well as good interactions between the PANI anchored graphene layer with the PVDF matrix. The dielectric permittivity of the composites of amino-­functionalized graphene (G-­NH2) and PVDF (PGN) shows the highest value of 59 for the PGN5 sample at 100 Hz. So it can be argued that the anchored conducting PANI chain from the GO surface has a strong influence for enormous increase of dielectric permittivity compared to that in their respective composites with PVDF. Figure 8.2d presents the variation of tan δ (dielectric loss) with frequency of the PVDF nanocomposites for varying concentration of G-­graft-­PANI over the frequency range 102–106 Hz. Generally, the dielectric loss of the materials originates from migration of space charge, direct current (dc) conduction, and movement of molecular dipoles. Dielectric loss decreases with increasing frequency and it is prominent for the PGP3 and PGP5 samples. The high dielectric loss values at the low frequency region are for the presence of the mobile charges within the polymer matrix and the lower value of dielectric loss at the higher frequency region is mainly due to polarization loss for dipolar and interfacial polarization. The nanocomposites exhibit increase in tan δ with increase in G-­graft-­PANI compared to the pure PVDF at the lower frequency (102 Hz) region due to dipole and interfacial polarization resulting in significant energy dissipation.18,19 The increased dispersion of G-­graft-­PANI in PVDF followed by increased interfacial adhesion in the nanocomposites, restricts the movement of the molecular dipoles resulting in an increase in dielectric permittivity and tan δ of the nanocomposite samples. The PGP5 composite shows a peak in the tan δ–frequency plot and it suggests a phase transition in the system but it is not prominent in other PGP samples because of lower filler concentration in the PVDF matrix.

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To comprehend the influence of β-­phase PVDF (determined from FTIR data) on the dielectric permittivity of the PGP and PGN composites dielectric permittivity is plotted with percentage β-­phase of PVDF in Figure 8.3a for both the samples at a constant frequency of 100 Hz. It is apparent from the plots that at lower β-­phase content (C=O groups of PMMA supramolecularly interact with the >CF2 groups of PVDF resulting in a homogenous dispersion of rGO functionalized PT-­g-­PMMA in the PVDF matrix. A solution mixing technique is used to make composite of the filler with PVDF, and FTIR and DSC data indicate that with increase in PT-­g-­PMMA functionalized rGO (PFG) polar β crystalline polymorph PVDF increases, producing about 90% β phase for 5 wt% filler in the composite. In accordance with the Maxwell–Wagner–Sillars interfacial polarization theory, high dielectric composite materials are formed by integrating conductive nanofillers in the insulating polymer matrix. The variation of dielectric permittivity (ε′) as a function of frequency (f) (Figure 8.9a) shows that on increasing the PFG concentration in the PVDF matrix the dielectric permittivity of PVDF (10.3) gradually increases, reaching a maximum value of 446.0 for 5 wt% filler content at 102 Hz. This gradual increase in dielectric permittivity with increase in PFG content is ascribed to the gradual increase of the microcapacitor network in the insulating PVDF matrix. The microcapacitors network pathway consists of high aspect ratio conducting PFG layers separated by a thin polar PVDF layer.43,44 As the PFG concentration in

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Figure 8.9  (a)  Dependence of dielectric permittivity with frequency at room temperature, (b) the log–log plot of the dielectric permittivity and frequency. Reproduced from ref. 42 with permission from Elsevier, Copyright 2016.

PVDF increases more and more, the polar β-­phase PVDF increases facilitating more microcapacitors network and conducting network formation, thus increasing dielectric constant showing a maximum for 5 wt% PFG content. It is important to note that with increasing frequency from 102 to 106 Hz the dielectric permittivity of PVDF composites decreases sharply and it is for the hindrance experienced by rotating CF2 dipoles causing a lag between the frequency of an oscillating dipole and that of the applied field. A power law of percolation theory is used to describe the relationship between the dielectric constant and filler volume fraction according to eqn (8.1), when the polymer functionalized graphene (PFG) approaches the percolation threshold.43,44 The best linear fitting of the log ε′ versus log(  fc − f) plots using fc = 0.4 (Figure 8.9b) of the power law [eqn (8.1)] yields s = 1.37. This low percolation threshold indicates a good dispersion of PFG in the matrix arising from good interfacial interactions between the PFG surfaces and PVDF matrix. In this system it is interesting to note that the conducting percolation threshold (vol% = 0.24, cf. Chapter 7) matches well with the dielectric percolation threshold (0.4 vol%) indicating that both the conductivity and dielectric properties originate from the same material structure, viz. the PFG network formation in the composite. Also the close match between the critical exponent values in conductivity and dielectric permittivity percolation plots supports the same origin of the two properties viz. the conducting network formation. A comparison with the covalently grafted G-­graft-­PANI/PVDF nanocomposite13 exhibiting a maximum ε′ value of 264 to that of the present noncovalently functionalized rGO/PT-­g-­PMMA/PVDF system (ε' = 446) for the same filler loading of 5 wt% at the same frequency of 100 Hz indicates the supremacy of the noncovalently functionalized system over the covalent functionalized graphene/PVDF system. This is because the aspect ratio and π-­conjugation in graphene, that attribute to the very high dielectric permittivity, remains

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unaltered in the noncovalently functionalized systems, but it decreases in covalent functionalized system despite the formation of the same amount of polar β-­phase PVDF (90%) in the composite.

8.3.1.2 Perylene Tetracarboxylic Acid (Py)/Exfoliated Graphene (EG)/PVDF Composite In order to further increase the dielectric properties of the technologically important polymer PVDF, recently Li et al.45 introduced a conductive nanofiller, perylene tetracarboxylic acid (Py), functionalized with electrochemically exfoliated graphene (EG) by π-­stacking interaction, into a PVDF matrix producing a percolative composite yielding very high dielectric constant with very low dielectric loss. The noncovalent functionalization between EG and Py by π-­stacking interaction not only maintains the high conjugation degree of EG, but also bestows EG with more carboxyl groups on its surface, facilitating strong dipolar interactions between the >CF2 group of PVDF and >C=O groups of PyEG along with hydrogen bonds between –COOH groups of PyEG and F atoms of PVDF. These specific interactions, together with the high surface area of EG, facilitate the transformation of the α-­phase to the polar β-­phase PVDF. The dielectric properties of PyEG/PVDF composites with different PyEG content are presented in Figure 8.10. The dielectric constants of PyEG/PVDF composites increase with the increasing PyEG content in PVDF and at 102 Hz the composite with 0.74 vol% of filler exhibits the highest dielectric constant value (∼1000) so far reported in the PVDF composites. This high increase is attributed to the increasingly large number of microcapacitors46 and the dielectric constant exhibits a common trend of a rapid decrease in dielectric constant with increasing frequency up to 104 Hz (Figure 8.10a), as the dipolar and interfacial polarization cannot match closer to the high applied frequencies.47 The gradual lowering of dielectric loss at frequency 104 Hz) may be attributed to the glass transition relaxation of the PVDF matrix, which is related to the micro-­Brownian motion (i.e. relaxation) of the whole PVDF chain.47,48 The inset of Figure 8.10a, shows the log–log plots of dielectric constant and frequency yielding percolation threshold value fc = 0.71% for PyEG/PVDF, and it is higher than that of EG/PVDF (  fc = 0.50%). This is because of the shielding effect of Py hinders the conjugation, hence more PyEG molecules are necessary than for EG to make a percolation network in the composites. In order to understand the advantages of PyEG nanosheets in improving the dielectric properties in the composites, the dielectric properties of PVDF composites containing different fillers are compared in Figure 8.10c and d. GO/PVDF composites show low dielectric constant and loss for poor conductivity of GO nanosheets.49 In the rGO/PVDF composite the dielectric constant is enhanced to a certain degree, but the high loss counterbalances the improvement in dielectric constant, and this loss is attributed to the aggregation of rGO due to worse interfacial interactions with the PVDF matrix.50

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Figure 8.10  (a)  Dielectric constant and (b) dielectric loss of PyEG/PVDF versus fre-

quency with different PyEG contents. (c) Comparison of the dielectric constants and (d) losses of GO/PVDF, rGO/PVDF, PyrGO/PVDF, EG/ PVDF and PyEG/PVDF composites with different filler content compiled at 1 kHZ. Reproduced from ref. 45 with permission from Elsevier, Copyright 2018.

On modification of rGO by Py molecules, the dispersity of PyrGO in the polymer matrix is improved; however, the dielectric constant still remains low due to the presence of a large number of defects on PyrGO sheets. The pristine EG nanosheets exhibit better distribution in the PVDF matrix than that of rGO nanosheets, but slight aggregations still remain, resulting in lower dielectric constant and higher loss.51 As presented in Figure 8.10c, PyEG at 0.74 vol% concentration in PVDF composite exhibits an amazing dielectric constant value of 480 at 1 kHz, and it is 55 times higher than that of neat PVDF at the same frequency. Moreover, compared with the same composition of EG/ PVDF, the dielectric losses are suppressed at a relatively low value of 0.27 (Figure 8.10d), The increased dielectric properties are attributed to the high conjugation degree and high interfacial interactions. To understand the underlying mechanism of the enhancement of dielectric properties of PyEG/PVDF compared to the dielectric properties of GO, rGO, Py modified rGO (PyrGO) and EG/PVDF composites, a rational model is presented in Figure 8.11 illuminating the synergistic effect of conjugation degree and interfacial interactions.

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Figure 8.11  Schematic  illustration of charge carrier localization and recombi-

nation as well as the interfacial interactions at the interfaces of the microcapacitor in (a) PyrGO/PVDF and (b) PyEG/PVDF composites. Reproduced from ref. 45 with permission from Elsevier, Copyright 2018.

From the view point of interfacial interactions, PyEG has higher interfacial interactions between PyEG and the PVDF matrix, as the former possesses a higher number of carboxyl groups yielding strong dipolar interactions between >C=O on PyEG and the >CF2 dipole of PVDF chains. Also the hydrogen bonds between carboxylic –OH and F atoms of PVDF lead to enhanced interfacial polarization at the interphase facilitating the conformational transformation from α-­ to polar β-­phase of PVDF.52 As a result microcapacitors at the heterointerface between PVDF matrix and PyEG are produced due to the formation of Schottky junctions facilitating charge injection and the charge carriers accumulate, surrounding the free space of PyEG nanosheets yielding a charged layer.53,54 Thus, the charge storage and loss behaviour can be explained by the charge transfer kinetics at the interfaces and the total effective charge increases with the width of accumulation layer. Due to higher conjugation of PyEG nanosheets from that of PyrGO larger free space permits more charge carriers to accumulate on their surfaces. However, in PyrGO nanosheets, due to the lack of sufficient free space for charge accumulation, recombination of localized carriers occurs, resulting in a high loss.

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So, with respect to graphene/polymer composites, the carrier recombination can be decreased by enhancing the effective conjugation degree of graphene to achieve lower loss. Thus the proposed model can reasonably explain the enhanced dielectric properties of these graphene based PVDF composites and this model may be used to optimize the dielectric properties for different composites with a variety of conductive fillers.

8.3.1.3 Poly(Sodium 4-­styrenesulfonate) Functionalized Graphene/Epoxy Nanocomposites Epoxy polymer is a thermoset resin generally used for fabricating versatile composites, adhesives and hardware components due to its excellent thermal, mechanical, chemical, dielectric and aging properties. To improve these properties Li et al.55 noncovalently functionalized rGO(g) with poly(sodium 4-­styrenesulfonate) (PSS) by π-­stacking interaction producing PSS-­g and successfully introduced it into the epoxy resin via in situ polymerization to form functional nanocomposites. To the PSS-­g dispersion in alcohol, epoxy resin is added and the mixture is ultrasonicated for several hours. On drying the solvent at 70 °C in vacuum, stoichiometric amounts of curing agent (DDM : epoxy = 1 : 4) and plasticizer (PPGDGE : epoxy = 1 : 4) are introduced into the mixture, stirred and, finally, the curing is done at 100 °C for 6 h in a vacuum oven. Different epoxy/PSS-­g composites with 0.5–2.0 wt% PSS-­g fillers are produced. In Figure 8.12 the frequency dependent electrical properties of epoxy/PSS-­g nanocomposites are presented. Here, both the conductivity and dielectric constant increase with increasing graphene content and the ac conductivity of nanocomposites with PSS-­g content 106 Hz). (Table 9.3) exhibits almost a constant value of 10.8 Ω for all the PAGD samples. In governing the PCE value of the cell the Rrc and Cµ values are important and depend upon the active material composition of the cell. The resistance (Rrc) and capacitance (Cµ) at the PAGD/electrolyte interface increase with enhancing GQD concentration, which is due to the poor conductivity and good charge stabilizing properties of the GQDs from those of PANI, respectively. Charges can accumulate well in the π-­stacked state of GQDs in the composites, and the poorer conductivity of GQDs compared with that of PANI arises for its shorter conjugation length. The high resistance (Rrc) and high lifetime values cause a decrease in back reaction of photo injected electrons which increase the short circuit current. This is true as Rrc and lifetime values increase, hence Jsc increases in the composites (cf. Table 9.3) with increasing GQD concentration, showing a maximum value for the PAGD3 sample which thus explains the highest PCE of 3.12%. The highest resistance (Rrc), capacitance (Cµ) and lifetime values in the PAGD3 sample may arise from its well organized rod-­like morphology produced from a highly ordered supramolecular structure at this optimum composition, yielding the highest PCE. Thus without using TiO2 as an active material we have been able to fabricate a DSSC with a PAGD hybrid showing a good PCE value.

9.3.5  R  eplacement of TiO2 Active Layer with Graphene/ Polymer Hybrid Xerogels Das et al.42 used xerogels of hybrid hydrogels containing GO, and a conducting polymer along with a gelator molecule to replace the TiO2. Dihybrid (GP) and trihybrid (GPPS) hydrogels are produced by using 5,5′-­(1,3,5 ,7-­tetraoxopyrrolo[3,4-­f]isoindole-­2,6-­diyl) diisophthalic acid (P), GO, and PEDOT:PSS, respectively. In the GP xerogel GO acts as acceptor and P acts as donor, whereas in the GPPS xerogel both the P and GO together act as joint acceptor and PEDOT:PSS acts as the donor for their relative electron accepting properties. The dc conductivity of GPPS xerogels are four to five

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orders higher compared to the GP and P xerogels due to the presence of the conducting polymer PEDOT:PSS. The GP and GPPS xerogels exhibit photocurrent on white light irradiation with 600% increase in photocurrent in the GPPS gel compared with that of GP gel. Also the on–off cycles display a stable photocurrent for the GPPS system. So DSSCs are fabricated with the GPPS xerogels where they act as active materials replacing TiO2 of DSSCs taking the N719 dye for varying PEDOT:PSS concentration (GPPS1, GPPS2, GPPS3 for 0.5, 0.71, 0.91% (w/v), respectively).42 The J–V characteristic curves of the hybrids (Figure 9.11a) indicate highest value of Jsc (10.2 mA cm−2), Voc (0.73 V) and fill factor (0.59) showing a maximum PCE of 4.46%. The incident photon to current efficiency (IPCE) curve of the system shows a wide absorption range (360–700 nm) with a maximum absorbance of ∼57%. The Nyquist plot of the DSSC obtained from impedance spectroscopy consists of three semicircles and the equivalent resistance–capacitance circuit yields the highest lifetime values of photoinjected electrons to be 3.2 ms explaining the highest PCE value of GPPS3 composite. The mechanism

Figure 9.11  (a)  Plot of J–V characteristics of GPPS1, GPPS2 and GPPS3 under AM

1.5 G light illumination of 100 mW cm−2. (b) Energy level diagram of the components of GPPS gels explaining the photocurrent behaviour of the FTO/GPPS gel/graphite DSSC with the N719 dye. Reproduced from ref. 42 with permission from the Royal Society of Chemistry.

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of photoconduction is presented in Figure 9.11b taking the LUMO energy levels of PEDOT:PSS (−2.4 eV),43 GO (−3.5 eV)44 and P (−4.25 eV).45 The N-­719 dye absorbs light promoting the electrons of the dye to their excited state and then enter into the LUMO of PEDOT:PSS and flow through the LUMO (conduction band) of GO and finally enter into the LUMO of P, leading to their flow in the external circuit (Figure 9.11b). The intimate mixing of the components (P, GO and PEDOT:PSS) of the xerogel assists the entire process of charge flow yielding a high value of PCE.

9.4  R  eplacement of the Pt Counter Electrode   with PFG Up to now, precious platinum (Pt) is preferred as counter electrode (CE) material for DSSCs for its good catalytic activity toward redox electrolyte along with good electrical conductivity and chemical stability. For commercial DSSC production the high amount of Pt material is a burden due to its price. So, efforts are now made on searching for alternative CE candidates and the fabrication of cost-­effective CEs is needed for developing low cost and advanced dye-­sensitized solar cells (DSSCs). Li et al.46 attempted the fabrication of conducting polymers such as polyaniline (PANi), polypyrrole (PPy), or poly(3,4-­ethylenedioxythiophene) (PEDOT) intercalated reduced graphene oxide (rGO) by mixing an aqueous solution of GO with a saturated aqueous solution of PPy, PANi or PEDOT followed by ultrasonication for 2 h to from a stable dispersion. The above mentioned dispersion is filtered and is moulded on titanium (Ti) foil and on a poly(ethylene terephthalate) (PET) substrate to form a film with an average thickness of ∼5 mm. After being rinsed with ethanol, the rGO/PPy, rGO/PANi and rGO/PEDOT CEs are dried and utilized as CEs for DSSCs. Figure 9.12a and b show the J–V characteristics curve of the DSSCs based on rGO/PPy, rGO/PANi and rGO/PEDOT on Ti foil or on a PET substrate, respectively.

Figure 9.12  Characteristic  J–V curves of DSSC fabricated with different CEs under one sun illumination on (a) Ti foil and (b) PET substrate. Reproduced from ref. 46 with permission from Elsevier, Copyright 2016.

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The DSSC with a Ti foil supported rGO/PPy CE yields a remarkable PCE = 6.23%, Jsc = 17.1 mA cm−2, Voc = 0.68 V and FF = 0.53 while the DSSC with a PET supported rGO/PPy CE yields PCE = 4.41%, Jsc = 14.2 mA cm−2, Voc = 0.589 V and FF = 0.527, which are much higher value than the photovoltaic parameters obtained with a pristine rGO CE based solar cell and that of other DSSCs. This is attributed to the rapid conversion from I3− to I−, accelerating the dye recovery facilitating constant electron photogeneration. It is necessary to mention that the determined cell efficiencies are higher than those from a pristine Pt based DSSC on a flexible PET substrate47 The DSSC with rGO/PPy, rGO/PANi and rGO/PEDOT electrodes display remarkable power conversion efficiencies of 6.23%, 5.73% and 4.65% for a Ti foil substrate, respectively, while the corresponding cell efficiencies are 4.41%, 2.93% and 2.05% on a PET substrate, respectively. The higher values for the Ti foil substrate over the PET substrate is due to its metallic nature facilitating the conduction of electrons in the outer circuit in addition to the polymer functionalized graphene. Thus, the photovoltaic performance is markedly enhanced by using conducting polymer intercalated rGO electrodes rather than a pure rGO based DSSC.

9.5  Improving Electrolyte Performance with PFG Researchers have not only used polymer functionalized graphene as photoactive and counter electrode materials but have also demonstrated it as an electrolyte material for DSSCs. Liquid electrolyte has low stability and limited robustness which limits the solar cell efficiency and stability. So it is very amazing to replace this liquid electrolyte containing I−/I3− in γ-­butyrolactone (GBL) with other materials without hampering the efficiency. To overcome this drawback, Nogueira and his group48 fabricated a reduced graphene oxide (rGO)-­polymer gel electrolyte as an alternative of liquid electrolyte. rGO is added to the copolymer poly(ethylene oxide-­co-­2-­(2-­methoxyethoxy)ethylglicidyl etherco-­allylglycidyl ether) [P(EO/EM/AGE)]), γ-­butyrolactone (GBL), LiI and I2, mixed well by stirring for 24 h and used as a polymer gel electrolyte. DSSCs are fabricated with the gel nanocomposite electrolytes containing different concentrations of rGO and, for comparison, a standard device was also fabricated using a commercial liquid electrolyte (EL-­HPE, Dyesol). It is clear that addition of rGO sheets in the gel electrolyte increases the short-­ circuit photocurrent (Jsc) and so also the Voc values increase from 0.60 to 0.65 V (Figure 9.13). The best PCE is achieved in the DSSC using 0.5 wt% of rGO in the gel electrolyte showing a value of 5.07% (Table 9.4). It is difficult to get electrolyte samples with a higher amount of rGO due to its low dispersion. The solar cell parameters indicate that rGO influences the performance of DSSC based on polymer gel electrolytes. On changing from classical polymer electrolytes to gel polymer electrolytes, more polyiodides species such as I3− and I5− are formed due to the increased mobility of ions present in the solution.49 These large species have low mobility, and tend to remain near the photoelectrode,

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Figure 9.13  Photocurrent–voltage  curves for the DSSCs using gel polymer electrolyte with different amounts of RGO compared with a DSSC using standard liquid electrolyte. The curves were obtained in AM 1.5 conditions (intensity of 100 mW cm−2). Reproduced from ref. 48 with permission from American Chemical Society, Copyright 2016.

Table 9.4  Electrical  parameters for the DSSCs using gel polymer electrolyte with different amounts of RGO compared to a DSSC using standard liquid electrolyte. Reproduced from ref. 48 with permission from American Chemical Society, Copyright 2016.

Electrolyte

Voc (V)

Isc (mA cm−2)

FF

η (%)

Gel/0.0% RGO Gel/0.1% RGO Gel/0.3% RGO Gel/0.5% RGO Liquid electrolyte

0.60 ± 0.08 0.64 ± 0.06 0.65 ± 0.05 0.65 ± 0.04 0.74 ± 0.02

12.63 ± 3.79 13.43 ± 3.22 13.90 ± 2.78 15.46 ± 3.31 15.43 ± 3.08

0.56 ± 0.10 0.55 ± 0.05 0.52 ± 0.05 0.51 ± 0.06 0.61 ± 0.05

3.99 ± 0.27 4.62 ± 0.79 4.59 ± 0.55 5.07 ± 0.97 6.97 ± 1.08

hence the incorporation of rGO in the gel polymer electrolytes acts in the opposite direction, increasing the values of Voc. This behaviour is related to the strong interaction of the polyiodide with the rGO sheets50 and allows the removal of polyiodide ions far from the photoelectrode interfaces, keeping them adsorbed on the surface of rGO sheets, recovering the Voc values. This result demonstrates the rGO-­polymer gel electrolyte as a new addition for electrolyte with low cost. Kowsari and Chirani51 also reported highly effective electrolyte additives using covalently functionalized GO with ammonium based ionic liquids for application in DSSC. GO is modified in two-­steps, first, with 1–6,

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hexamethylenediamine and finally, reacting with alkyl-­halides. The covalently bonded GO-­hexamethylene tributyl-­ammonium iodide (GO-­HMA-­ TBAI) and GO-­hexamethylene trimethyl-­ammonium iodide (GO-­HMA-­TMAI) are used as additives to the liquid electrolyte containing I−/I3−. The addition of ionic liquids minimizes the aggregation occurring in DSSC and improves the performance of DSSCs. The effects of different co-­additives on the photovoltaic performance of the DSSCs are shown in Figure 9.14. It is evident that DSSC fabricated with bare GO-­additive electrolyte (E1, consisting of I2, LiI, TBP in an ILs solvent) as a standard electrolyte demonstrates a low conversion efficiency of 3.96% with a low short circuit current density (Jsc) of 7.141 mA cm−2 and an open circuit voltage (Voc) of 0.73 V. The figure clearly demonstrates a strikingly enhanced Jsc and PCE values with the ionic liquid additive containing composite electrolyte DSSC fabricated with E1 + 0.4 wt% of GO (E2), shows Jsc of 9.175 mA cm−2 and Voc of 0.74 V yielding a PCE of 5.09. The DSSC based on E3 electrolyte (E1+ 4 wt% of GO-­HMA-­ TPMI) produces Jsc of 13.11 mA cm−2, Voc 0.75 V and PCE = 6.78%; while the Jsc and PCE of a DSSC based on electrolyte E4 (E1+ 4 wt% of GO-­HMA-­TBAI) shows a maximum Jsc of 16.847 mA cm−2 and PCE = 8.33%. It is evident that the difference between the values of Jsc obtained using the electrolyte with GO-­HMA-­TBAI as additive and electrolyte without any co-­additive is more than two times. From this study it is apparent that the IL-­functionalized GO forms a molecular bridge for movement of electrons in the ionic liquid based electrolyte causing a significant rise in Jsc from 7.141 to 16.847 mA cm−2. At an optimized additive/electrolyte ratio of 4 wt%, the PCE of DSSCs increases

Figure 9.14  Current  density–voltage (J–V) characteristics of DSSCs based on different electrolytes: E1 (I2, LiI, TBP in an IL solvent containing PMII and DMII), E2 (E1 + 4 wt% of GO), E3 (E1 + 4 wt% of GO-­HMA-­TMAI), and E4 (E1 + 4 wt% of GO-­HMA-­TBAI). Reproduced from ref. 51 with permission from Elsevier, Copyright 2017.

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amazingly to 8.33% from 3.96% of the referenced ionic liquids based device. The resulting DSSCs have exhibited better electron diffusion of I− and I3− and ionic conductivity which are the causes of improvement of Jsc. So this work would provide an effective way to develop graphene nanocomposite as a great potential for practical application. So polymer functionalized graphene (PFG) is very much effective in fabricating high performance DSSC without using the costly TiO2 active material of photo-­electrode, Pt -­ counter electrode also improves the electrolyte performance of DSSCs.

9.6  Perovskite Solar Cell The discovery of the perovskite solar cell attracted a huge amount attention because of its high power conversion efficiency.52–54 The perovskite solar cell consists of an electron transporting layer (ETL, n-­t ype semiconductor), a hole transporting layer (HTL, p-­t ype semiconductor) and an organic–inorganic halide perovskite, which is an intrinsic semiconductor (i). This perovskite also acts as a light absorber generating holes and electrons, which flow through the HTL and ETL, respectively, to the corresponding electrodes generating current on illumination with light. These organic–inorganic halide perovskite solar cells are considered to be the most promising material in the field of photovoltaics because of their ease of fabrication with the use of inexpensive materials yielding high PCE values.55 The perovskite solar cells (PSCs) have two types of cell architectures, e.g. p–i–n or the inverted n–i–p configuration. The most important drawback of PSCs is their low stability, so for commercialization of perovskite solar cells researchers are aiming at producing not only more high cell efficiency but are also targeting increasing their environmental stability. Here we shall discuss only the role of polymer functionalized graphene (PFG) in the development of HTL and ETL.

9.6.1  P  oly[(5,6-­difluoro-­2,1,3-­benzothiadiazol-­4,7-­diyl)-  ­alt-­(3,3‴-­di(2-­octyldo decyl)-­2,2;5,2;5,2″-­quaterthiophen-­ 5,5‴-­diyl)] as Hole Transporting Layer (HTL) Due to low electron mobility in TiO2 and low stability of ZnO, a graphene–polymer composite is introduced as an alternative for the electron transporting layer (ETL) in perovskite solar cells (PSCs). Loh and his group demonstrated the use of a mesoporous graphene/polyaniline composite as an ETL for PSCs.56 In consideration of low-­temperature processing requirements and the long-­ term stability of PSCs,57,58 they reported a low-­temperature processed mesoporous graphene/polymer (mp–GP) ETL for making high performance PSCs. The network structures of conducting graphene and granular-­like polyaniline having pores serve as fast electronic channels for the infiltration of active layers, causing complete surface coverage of perovskite. Compared to pure

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graphene, the mp–GP ETL exhibits an efficiency enhancement and also protects the perovskite layer due to its chemical inertness for its ability to encapsulate the perovskite crystals, protecting them from moisture and reactive interface formation. Figure 9.15 presents the device structure, J–V plot and energy diagram of an mp–GP based PSC. In the conventional device structure, the perovskite is sandwiched between an ETL and an HTL with low LUMO and high HOMO levels, respectively, to achieve high performance PSCs. Here poly[(5,6-­difluoro-­ 2,1,3-­benzothiadiazol-­4,7-­diyl)-­alt-­(3,3‴-­di(2-­octyldo decyl)-­2,2′;5′,2″;5″,2‴-­ quaterthiophen-­5,5‴-­diyl)] (PffBT4T-­2OD) is used as the HTL, and compared to conventional HTL spiro-­OMeTAD, PffBT4T-­2OD has a deeper HOMO which provides a larger value of open-­circuit voltage (Voc) in the PSC.59 The large LUMO offset between the PffBT4T-­2OD/perovskite interface is constructive to hinder the electrons from transporting from the perovskite to the Ag anode. The work function of the mp–GP film (4.4 eV) causes a high electron extraction barrier of 0.47 eV at the mp–GP/perovskite interface. Appropriate interfacial engineering with low work function (2.3 eV) Cs2CO3 60 is applied on the ETL to make a high performance photovoltaic cell. The tuneable energy level of the mp–GP composite with Cs2CO3 is pertinent to reduce the interfacial energy barrier and higher electron mobility

Figure 9.15  (a)  Schematic view of the device structure. (b) J–V characteristics of PSCs with ITO/mp–GP or rGO/Cs2CO3/CH3NH3PbI3/PffBT4T-­2OD/Ag. The performance of devices without Cs2CO3 are also added for comparison. (c) Energy level diagrams of the device. Reproduced from ref. 56 with permission from American Chemical Society, Copyright 2016.

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2

−1

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of the mp–GP (1.7 × 10 cm V s ) over conventional ETL such as TiO2 (1.7 × 10−4 cm2 V−1 s−1) suggests that mp–GP can be a good alternative ETL in PSCs. Figure 9.15b shows the J–V characteristic plots of PSCs using different ETLs: mp–GP, rGO, mp–GP-­Cs2CO3, and rGO-­Cs2CO3. All ETLs are processed under mild (